Mechanism of Irradiation Assisted Cracking of Core ... · Table 3: Summary of grain boundary composition measurements* on post-irradiation annealed CP-304 (irradiated to1.0 and 2.5

DE-FG07-99ID13768 Final Report

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Nuclear Engineering Education Research (NEER)Final Scientific/Technical Report

Mechanism of Irradiation Assisted Cracking of Core Components in LightWater Reactors

Grant #DE-FG07-99ID13768Project No. 99-1637

Submitted by:University of Michigan

Gary S. Was, PIMichael Atzmon, co-PI

Lumin Wang, co-PIJeremy Busby

Submitted on:

April 30, 2003


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Executive Summary

The overall goal of the project is to determine the mechanism of irradiation assisted stress

corrosion cracking (IASCC). IASCC has been linked to hardening, microstructural and

microchemical changes during irradiation. Unfortunately, all of these changes occur

simultaneously and at similar rates during irradiation, making attribution of IASCC to any one of

these features nearly impossible to determine. The strategy set forth in this project is to develop

means to separate microstructural from microchemical changes to evaluate each separately for

their effect on IASCC. In the first part, post irradiation annealing (PIA) treatments are used to

anneal the irradiated microstructure, leaving only radiation induced segregation (RIS) for

evaluation for its contribution to IASCC. The second part of the strategy is to use low

temperature irradiation to produce a radiation damage dislocation loop microstructure without

radiation induced segregation in order to evaluate the effect of the dislocation microstructure

alone.

A radiation annealing model was developed based on the elimination of dislocation loops

by vacancy absorption. Results showed that there were indeed, time-temperature annealing

combinations that leave the radiation induced segregation profile largely unaltered while the

dislocation microstructure is significantly reduced. Proton irradiation of 304 stainless steel

irradiated with 3.2 MeV protons to 1.0 or 2.5 dpa resulted in grain boundary depletion of

chromium and enrichment of nickel and a radiation damaged microstructure. Post irradiation

annealing at temperatures of 500 – 600°C for times of up to 45 min. removed the dislocation

microstructure to a greater degree with increasing temperatures, or times at temperature, while

leaving the radiation induced segregation profile relatively unaltered. Constant extension rate

tensile (CERT) experiments in 288°C water containing 2 ppm O2 and with a conductivity of 0.2

mS/cm and at a strain rate of 3 x 10-7 s-1 showed that the IASCC susceptibility, as measured by

the crack length per unit strain, decreased with very short anneals and was almost completely

removed by an anneal at 500°C for 45 min. This annealing treatment removed about 15% of the

dislocation microstructure and the irradiation hardening, but did not affect the grain boundary

chromium depletion or nickel segregation, nor did it affect the grain boundary content of other

minor impurities. These results indicate that RIS is not the sole controlling feature of IASCC in

irradiated stainless steels in normal water chemistry.


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The isolation of the irradiated microstructure was approached using low temperature

irradiation or combinations of low and high temperature irradiations to achieve a stable,

irradiated microstructure without RIS. Experiments were successful in achieving a high degree

of irradiation hardening without any evidence of RIS of either major or minor elements. The low

temperature irradiations to doses up to 0.3 dpa at T<75°C were also very successful in producing

hardening to levels considerably above that for irradiations conducted under nominal conditions

of 1 dpa at 360°C. However, the microstructure consisted of an extremely fine dispersion of

defect clusters of sizes that are not resolvable by either transmission electron microscopy (TEM)

or small angle x-ray scattering (SAXS). The microstructure was not stable at the 288°C IASCC

test temperature and resulted in rapid reduction of hardening and presumably, annealing of the

defect clusters at this temperature as well. Nevertheless, the annealing studies showed that

treatments that resulted in significant decreases in the hardening produced small changes in the

dislocation microstructure that were confined to the elimination of the finest of loops (~1 nm).

These results substantiate the importance of the very fine defect microstructure in the IASCC

process.

The results of this program provide the first definitive evidence that RIS is not the sole

controlling factor in the irradiation assisted stress corrosion cracking of austenitic stainless steels

in normal water chemistry. Earlier research has suggested that RIS may not play the dominant

role it has recently been afforded, but the results of this program are the first to definitively

establish this role. The program has also shown that the fine defect structure is implicated in the

IASCC process and likely plays a role in the observed cracking of core components. The results

of this project provide the basis and the motivation for further investigation into the mechanism

of IASCC in which fine defect clusters may play a crucial, if not a defining role.


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TABLE OF CONTENTS

Title Page 1Executive Summary 2Table of Contents 4List of Tables 5List of Figures 61.0 Introduction 81.0 Background 9

1.1 Possible Mechanisms of IASCC 91.1 Objective 121.1 Approach 121.1 The Research Plan 13

1.0 Experiment 161.1 Material Preparation 161.1 Proton Irradiation 161.1 Radiation Induced Segregation 171.1 Microstructure Analysis 181.1 Post Irradiation Hardness Measurement 181.1 Stress Corrosion Cracking Experiments 19

1.0 Results and Discussion 201.1 Post-Irradiation Annealing: Simulation Results 20

1.1.1 Simulated annealing of dislocation loops and hardening 211.1.1 Thermodynamic considerations 231.1.1 Kinetic considerations 25

1.1 Post-Irradiation Annealing: Experimental Results 281.1.1 RIS 281.1.1 Dislocation loops 351.1.1 Hardness 401.1.1 IASCC of post-irradiation annealed stainless steel 43

1.1 Discussion of Post-Irradiation Annealing Results 461.1.1 Separation of RIS and Loops 461.1.1 RIS of Cr and Ni and IASCC 471.1.1 RIS of Si and P and IASCC 481.1.1 Dislocation loops and IASCC 491.1.1 Hardening and IASCC 501.1.1 Other potential contributors to IASCC 501.1.1 Other elements 511.1.1 Combination of RIS and microstructure 511.1.1 Small defect clusters 521.1.1 Implications for IASCC 55

4.4 Isolating the Dislocation Microstructure 575.0 Conclusions 666.0 References 687.0 Publications 718.0 Students 72


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List of Tables

Table 1: Bulk composition of HP-304L and CP-304 alloy as determined by electron microprobeanalysis (wt% and at%)

Table 2: Summary of grain boundary composition measurements on post-irradiation annealedHP-304L (irradiated to1.0 at 360°C). All results are listed in at%.

Table 3: Summary of grain boundary composition measurements* on post-irradiationannealed CP-304 (irradiated to1.0 and 2.5 dpa at 360°C). All results are listed in at%.

Table 4: Summary of dislocation loop analysis on post-irradiation annealed HP-304L and CP-304 (irradiated to 1.0 and 2.5 dpa at 360°C).

Table 5: Summary of hardness analysis on post-irradiation annealed CP-304 (irradiated to 1.0and 2.5 dpa at 360°C).

Table 6: Summary of CERT test results performed on post-irradiation annealed CP-304 samples(1.0 and 2.5 dpa).

Table 7: Hardness of CP304 SS following irradiation and annealing.

Table 8. Hardness before and after various irradiations to 0.5 dpa at either T<75°C or at 360°C.All hardness units are in kg/mm2.

Table 9. Results of constant extension rate test in 288°C BWR normal water chemistry.


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List of Figures

Figure 1. Schematic of the development of dislocation loop microstructure, hardening and RISalong with the increase in IASCC with irradiation dose.

Figure 2. Experimental plan for isolating microstructure and microchemistry in irradiated304SS.

Figure 3. Displacement rate profile for 3.2 MeV protons in stainless steel as calculated by theMonte Carlo program TRIM 90.

Figure 4: Comparison of simulated annealing of fraction of as-irradiated grain boundary Crdepletion and total loop line length as a function of time for anneals at 400°C, 500°C, and600°C. Simulation for CP-304 irradiated to 1.0 dpa at 360°C.

Figure 5: Comparison of activation energies for removal of dislocation loops and RIS duringsimulated post-irradiation anneal of HP-304L irradiated to 1.0 dpa at 360°C.

Figure 6: Simulation of annealing of Cr segregation as a function of time during post-irradiationannealing at 500°C. Simulation with as-irradiated dislocation density and no dislocation density.

Figure 7: Cr segregation profiles for HP-304L and CP-304 irradiated with 3.2 MeV protons at360°C to 1.0 dpa and 2.5 dpa and post-irradiation annealed. The as-irradiated profile (opensymbols) is shown in each figure. The 0 nm position is the grain boundary for all profiles.

Figure 8: Annealing of Cr segregation as a function of Fe-diffusion distance. The % of as-irradiated minimum measured Cr is plotted for all conditions. Data points for CP-304 at 1.0 dpaand 2.5 dpa have been shifted left and right, respectively, for clarity.

Figure 9: Comparison of measured annealing of Cr segregation with simulated annealing usingthe MIK model.

Figure 10: Bright field images of dislocation loop populations in HP-304L and CP-304irradiated with 3.2 MeV protons at 360°C to 1.0 dpa and 2.5 dpa and post-irradiation annealed.

Figure 11: Annealing of dislocation microstructure as a function of Fe-diffusion distance. Thefraction of the as-irradiated loop line length associated with the dislocation population is plotted.

Figure 12: Comparison of simulated and measured annealing of dislocation loop line length.

Figure 13: Annealing of measured hardness as a function of Fe-diffusion distance. The fractionof the as-irradiated change in yield stress calculated from hardness measurements is plotted.

Figure 14: Comparison of measured total crack length or %IG remaining as a function and Fe-diffusion distance. Also, the irradiated surface of CP-304-2.5 dpa sample post-irradiationannealed at 450°C and 500°C for 45 min are shown. Samples strained at 3 x 10-7 s-1 in water at


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288°C, 0.2 mS/cm, and 2 ppm O2. A line has been added to the fracture surface of the sampleannealed at 500°C for 45 min. to indicate the irradiated surface of the sample.

Figure 15: Comparison of annealing behavior of major alloying elements, a.), minor alloyingelements, b.), dislocation loop line length, c.) and hardness, d.) with cracking susceptibility forCP-304 irradiated to 1.0 and 2.5 dpa.

Figure 16: Schematic of annealing behavior of mixed population of black-dot damage anddislocations. During annealing, small defect clusters are removed or dissociate into individualinterstitials which may then be absorbed by dislocation loops. This process results in apopulation of larger loops with no change in density.

Figure 17: Comparison of measured and simulated annealing of small dislocation looppopulations from proton irradiations to 1.0 dpa at 360°C (plotted as Nd in c.)).

Figure 18: Comparison of slip bands in CP 304 SS irradiated to 1.0 dpa at 360°C and annealedat a.) 450°C/45 minutes (IG with 65 slip band systems/mm2) or b.) 500°C/45 minutes (no IG and25 slip band systems/mm2). Also shown in c.) are the slip bands observed on a CP 304 SSsample irradiated to 0.3 dpa at 360°C (no IG with 22 slip band systems/mm2) and d.) 0.3 dpa at <75°C (no IG with 28 slip band systems/mm2). Samples strained at 3.5 x 10-7 s-1 in water at 288°C,0.2 mS/cm, and 2 ppm O2.

Figure 19. Change in hardness of CP304 SS following annealing of samples irradiated to 0.3dpa at T<75°C.

Figure 20. Composition profiles for Cr, Ni, Fe and Si across grain boundaries of CP304 SSfollowing low temperature irradiation at T<75°C showing the lack of RIS.

Figure 21. Scheme for achieving a stable irradiated microstructure by combining hightemperature (T=360°C) and low temperature (T<75°C) irradiation in increments of 0.25 dpa.

Figure 22. Stress-strain curves of CP304SS samples tested in normal water chemistry consistingof 288°C water containing 2 ppm O2 and conductivity of 0.2 mS/cm. Samples were irradiated tothe following conditions: a) 0.25 dpa @ T<75°C + 0.25 dpa @ 360°C, b) 0.25 dpa @ 360°C +0.25 dpa @ 360°C, c) 0.25 dpa @ 360°C + 0.25 dpa @ T<75°C, and d) 0.25 dpa @ 360°C +0.25 dpa @ T<75°C + 0.25 dpa @ 360°C.

Figure 23. Change in hardness of CP304SS following irradiation to 0.3 dpa and annealing at288°C for various times up to 24 hours.

Figure 24: Comparison of annealing of samples irradiated at low temperatures with modelsimulations of annealing of small defect clusters.

Figure 25: Comparison of annealing of samples irradiated at low temperatures with predictionsfrom model developed by Simonen.


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1.0 INTRODUTION

Irradiation assisted stress corrosion cracking (IASCC) describes the intergranular

cracking that occurs in irradiated structural components in water reactors. The general feature

common to these failures is increased susceptibility of various non-sensitized austenitic stainless

steels (SS) with neutron fluence. Incidents of IASCC were first reported in the early 1960’s [1]

in 304, 304L and 347 SS fuel rod cladding in BWRs, [2-4] and more recently in 304 SS fuel

cladding in the Connecticut Yankee reactor (PWR). [5] IASCC also occurs in 304 SS control

rod absorber tubes, fuel bundle cap screws, control rod blade handles, sheaths and follower

rivets, plate type control blades and instrument dry tubes in BWRs. Brown and Gordon [6,7]

reported detailed IASCC field histories in BWRs, including cracking in Alloy 600 shroud head

bolts (first observed in 1986), stainless steel safe ends (1984) and in-core instrumentation tubes

(1984). At the West Milton PWR test loop, intergranular failure of vacuum annealed type 304

stainless steel fuel cladding was observed [8] in 316°C ammoniated water (pH 10) when the

cladding was stressed above yield. Similarly, IASCC was observed in creviced stainless steel

fuel element ferrules in the Winfrith SGHWR, [9] a 100 MWe plant in which light water is

boiled within pressure tubes.

IASCC occurs irrespective of reactor type. Specific BWR vs. PWR comparisons were

performed using in-core swelling tubes [1,10] fabricated from a variety of commercial and high

purity heats of types 304, 316 and 348 stainless steel and Alloys X-750, 718 and 625. Based on

identical strings of specimens placed in fuel rod locations, there was little distinction in the

IASCC response between the two reactor types. [11] Thus, it is becoming increasingly evident

that the problem is widespread without regard to environment or alloy, and that numerous core

components may be susceptible to this form of degradation. A recent review by Andresen [1]

identified failures in some 17 components, spanning 6 iron- or nickel-base alloys and 4 reactor

designs. Given that a threshold in fluence has been identified, many of the susceptible

components may just now be coming to our attention and the extent of the problem may continue

to expand.

The future of current light water reactors and future water reactor concepts may well rest

on the solution to the IASCC problem. It is clear that stress, an irradiated microstructure and an


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aggressive environment are all required for IASCC to occur. Further, the microstructural

features due to irradiation are generally well known. What is not known is which of these

features is responsible for IASCC and why. This program seeks to answer the first part of this

question; what microstructural features are responsible for IASCC. The following section

provides background on the IASCC phenomena along with the specific objective of this program

and the approach to the problem.

2.0 BACKGROUND

2.1 Possible Mechanisms of IASCC

Since irradiation by high energy neutrons can affect both the material properties and the

environmental chemistry, IASCC is a complex problem which is neither easy to investigate

practically nor simple to understand mechanistically. [10] The mechanisms which have been

proposed to explain IASCC are categorized into either effects of radiation on the “environment”,

such as corrosion potential and ion chemistry, or “persistent” effects on the microstructure such

as radiation hardening and creep, and on the microchemistry such as radiation-induced

segregation of alloying elements and impurities. While the environment can affect the degree of

cracking, it is the “persistent” effect of irradiation on the material that is primarily responsible for

its susceptibility. This point is supported by the observation of a distinct threshold fluence (1

dpa in BWRs and 3-5 dpa in PWRs) at which IASCC occurs. [10] The existence of a threshold

fluence for IASCC both in-situ and in post-irradiation tests indicates that "persistent" radiation

effects (microstructural and microchemical changes) are crucial for IASCC. This position was

emphasized by a jointly sponsored (DOE & EPRI) Radiation Assistance Task Force on radiation

materials science at a workshop held in March, 1998. [12] Thus, the key to understanding the

mechanism is understanding the radiation effect on the material. This amounts to two possible

irradiation-induced changes; those that alter the microstructure and those, which alter the

microchemistry.

Microstructure effects are predominantly faulted dislocation loops and an increase in the

dislocation network density. Radiation induced precipitation (phase changes) and bubble

formation do not occur at typical LWR temperatures and doses and voids are rarely observed.


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The dislocation loops range from small, unresolvable “black dots” to larger Frank-type faulted

loops between 8 and 12 nm in diameter. From the standpoint of cracking, their significance is

twofold; they result in lattice hardening by acting as hard barriers to dislocation motion, and they

cause strain localization (dislocation channeling), in which plastic deformation occurs

inhomogeneously in groups of slip bands through which a lead dislocation has passed and

removed the obstacles for subsequent dislocation motion. While it is unknown how such

localization of strain relates to IG cracking, it remains a strong possibility that there is a relation

between the deformation mode and cracking. However, experiments conducted to date are

inconclusive due to the fact that strain localization and the causative microstructure always occur

simultaneously with a change in microchemistry at the grain boundaries.

Local composition changes occur as a result of radiation induced segregation (RIS) of

both major alloying elements and impurities. In austenitic alloys, these changes can be quite

substantial with the contents of majority elements (Cr, Ni) changing by a factor of 2 or more and

impurities enriching by factors approaching 100. There is strong suspicion that such significant

composition changes at the grain boundary are responsible for the observed IASCC. In fact, a

very likely mechanism for IASCC observed both in-plant and in the laboratory is chromium

depletion due to RIS. The loss of a protective chromium film due to depletion of chromium at

the grain boundaries resulting from chromium carbide precipitation during heat treating was the

cause of the BWR pipe cracking problem in the 70s and 80s. [13] In fact, high temperature,

oxidizing environments have been shown repeatedly to produce intergranular stress corrosion

cracking in chromium depleted austenitic iron- and nickel-base alloys. [14] Although the

chromium-depleted zones caused by RIS are much narrower (by 100x) than those produced by

carbide precipitation, they have also been associated with IGSCC in oxidizing media, such as

exist in a BWR. However, the problem is that IASCC also occurs in PWRs, which have

reducing environments that are known to be benign to chromium depletion. Herein lies the

major problem: while cracking morphology and dose dependence are nearly identical in

oxidizing (BWR) and reducing (PWR) systems (except for a higher dose threshold in PWRs),

chromium depletion can explain cracking in the former but not in the latter. Further, it is


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unlikely that impurity segregation is the determining factor since cracking susceptibility is no

different for high-purity heats both in-plant and in the laboratory. [15]

Figure 1. Schematic of the development of dislocation loop microstructure, hardening and RIS alongwith the increase in IASCC with irradiation dose.

Significant amounts of cracking data can be shown to correlate to both the development

of the dislocation microstructure and radiation induced segregation. Figure 1 schematically

depicts the problem in attributing IASCC to a particular irradiation induced microstructural

change. As the figure shows, cracking increases with dose, but so does the dislocation

microstructure, hardening (which results from the microstructure) and the degree of radiation

induced segregation. Hence, while it has been shown that changes to the material control

IASCC, it remains unclear whether microstructure changes or microchemistry changes or some

combination of the two are controlling. This point was underscored by the Cooperative IASCC

Research Program (CIR), an international group of reactor vendors, utilities and regulators,

which identified the uncertainty of the relative roles of microstructure and microchemistry as the

principal barrier to solving the IASCC problem. It is possible that some other factor may be

controlling cracking. However, the search for an unknown is pointless until it can be shown that

neither microstructure nor microchemistry control IASCC. The solution to the IASCC problem


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must begin with an understanding how each of these principal irradiation-induced changes affect

cracking in oxidizing and reducing environments, as stated in the following objective.

2.2 Objective

The objective of this program is to determine the mechanism of IASCC of core

components in austenitic light-water reactor core components. The premise is that IASCC is

caused either by radiation induced segregation (RIS) at the grain boundary, or by the radiation

induced dislocation microstructure. This program will week to identify the causative factors of

IASCC by first isolating, and then independently evaluating the roles of microstructure and

microchemistry in SCC susceptibility.

2.3 Approach

We will identify the “persistent” effect responsible for IASCC by separating the

microstructure changes from the microchemistry changes. That is, we will create a

microstructure typical of that resulting from neutron irradiation in reactor to several dpa, but with

little change in microchemistry from the unirradiated state. Similarly, we will create a

microchemistry typical of that resulting from neutron irradiation in-reactor to several dpa, but

with little change in microstructure from the unirradiated state. In this way, we can perform

SCC tests to isolate the key material condition that is responsible for the observed cracking. We

do not expect to completely eliminate one feature while completely retaining the other, but we do

expect that the feature of interest will strongly dominate the other. We will then be able to

conduct more specific investigations into the mechanism by which microstructure or

microchemistry changes affect the IG cracking process.

We propose to create these microstructures using proton irradiation. This has developed,

over the past 10 years, into a proven technique for producing representative microstructures in

both 304 and 316 SS and in high purity analogues without significant residual radioactivity, at a

much lower cost and without the large amounts of time and expense required to conduct in-

reactor irradiations. In fact, this is the only practical way to conduct a mechanistic study of the


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problem of IASCC.The end result of this program will be the identification of the material changes that affect

IASCC and a better understanding of the mechanism. Until such changes are identified, no further

progress can be made on identifying the mechanism and solving the problem. Solving the IASCCproblem is a long-term goal, but not the objective of this proposal. An understanding of the mechanism

will allow for the development of mitigation strategies for existing core components and also the

development of radiation-resistant alloys or microstructures for replacement components and for

advanced reactor designs. These developments are essential to the extension of existing plants to longer

lives and the prevention of IASCC in advanced reactor designs. This strategy has the support of EPRIas evidenced by the substantial cost sharing provided in the budget.

2.4 The Research Plan

The plan diagrammed in Fig. 2 is based on isolating the relevant LWR microstructure and

microchemistry. Experiments will be conducted on a high purity Fe-Cr-Ni-Mn alloy (304 SS

analogue) and selected experiments will be conducted on archive samples of a commercial 304

SS alloy. The high purity alloy (selected to avoid complications from impurities) will be cold-

worked and annealed to produce a grain size of 10-12 micrometers. The commercial alloy was

irradiated in reactor and was shown to be susceptible to IASCC via constant extension rate

tensile tests. [15] An extensive characterization of its microstructure and microchemistry has

been conducted by PNNL and we have shown that both can be replicated by proton irradiation.

[17-21] Hence, this alloy represents the ideal case for studying microstructure and

microchemistry effects using proton irradiation. Irradiation will be conducted using 3.2 MeV

protons generated in the Tandetron accelerator at the Michigan Ion Beam Laboratory at the

University of Michigan. This produces a nearly flat damage profile over the first 35 mm,

followed by a damage peak at ~40 mm. [22] Characterization of microstructure and

microchemistry changes will be made in the flat damage region, away from the surface and the

damage peak. Microstructure characterization will be made in TEM and microchemistry

characterization will be made using STEM-EDS.!


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14Figure 2. Experimental plan for isolating microstructure and microchemistry in irradiated 304SS.

IASCC experiments will be conducted in a specially designed flowing autoclave system in which

water chemistry (corrosion potential, oxygen content, solution conductivity, pH, flow rate),

temperature, pressure and strain rate can be controlled on four samples tested independently. [23]

Cracking susceptibility is evaluated by the number and length of cracks in the irradiated region

of the sample over the course of a constant extension rate test. (Note that because the irradiatedlayer is 40 mm deep and the bulk is ductile, the sample will not readily fail upon formation of the

first crack.) We will also conduct a limited number of interrupted tests to determine crack

nucleation time and an estimate of the crack growth rate. [24,25]It is important to note that this program involves multiple experimental and analysis

capabilities, each of which may constitute a project in itself. We are in a position to bring all of

these techniques to bear on this problem in a single proposal because of the extensive effort that

has been conducted over the past 10 years to develop and validate this technique for studyingneutron irradiation damage. Thus, the specifics of proton irradiation (including dose and

MicrostructureMicrostructureIsolationIsolation

1.) L.T. irradiation + He implantation + annealto form radiation microstructure w/o Cr depletion

2.) Irradiate pre-enriched grain boundaries to remove pre-enrichment and introduce damage

Irradiation at 360°C + annealto form Cr depleted grain boundaries

No No YesYes

indicates• microchemistry or• combination

indicatesmicrochemistry

indicatesmicrostructure

check hardening vsstrain localization

create hardening via cold-working

No Yesmust bestrain

localizationmust be

hardening

No Yes

microchemistryother than Cr

depletionis involved

must beCr depletion

IGSCCsusceptible?

IGSCC*susceptible?

IGSCC inreduc. env?

IGSCC*susceptible?

*IASCC experiments are conducted in oxidizing environments (288°C, pH 7.0, O2~2 ppm, Ecorr ~ +150 mVSHE, flowing system)

and at a constant extension rate of 1 x 10-7 s-1.

indicates• microstructure or• combination

MicrochemistryMicrochemistryIsolationIsolation


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temperature control), microstructure analysis, microchemistry analysis and SCC testing in

relevant BWR and PWR environments have all been established. [26] The verification of the

method has also been well established. [27] This program will benefit from prior work in thatthe entire project can be devoted to solving the IASCC problem, rather than to developing a

capability.


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3.0. EXPERIMENT

3.1 Material Preparation

Two 304 alloys, a high purity, low carbon alloy made by General Electric, Schenectady,

and a commercial purity 304 stainless steel alloy from ABB Atom, with the compositions listed

in Table 1 (in at%) were used in this study. The UHP-304L, in the form of 12 mm thickness

bars, were solution annealed at 1100°C for 1 hour to produce a homogeneous microstructure and

microchemistry, and a grain size of approximately 100-200 mm. The alloys were then cold

rolled from 12 mm down to 2 mm to provide a reduced grain size upon recrystallization. These

bar samples were then wet-polished using 320-600 grit silicon carbide (SiC) paper to remove the

mechanical damage introduced during machining. The final annealing treatment in a flowing

argon atmosphere at 850 °C for 30 min recrystallized the alloys to achieve grain sizes of

approximately 10 mm. The commercial alloy was used in the as-received condition and had a

grain size of approximately 50 microns. The CP304 alloy was archive material from the same

heat that had also been irradiated in the Bärseback 1 reactor in Sweden. [28] Microstructural,

microchemical, hardness and IASCC data was measured previously as a function of dose. [29]

Table 1: Bulk composition of HP-304L and CP-304 alloy as determined by electron microprobe analysis(wt% and at%)

Cr Ni Fe Mn Mo Si B C N P SHP-304L

wt% 19.7 9.5 69.7 1.12 0.02 0.01 -- 0.006 <0.001 0.001 0.002at% 20.9 9.0 69.0 1.1 0.01 0.02 -- 0.028 <0.004 0.002 0.003

CP-304

wt% 18.3 8.5 70.6 1.38 0.37 0.65 <0.0004 0.035 0.068 0.03 0.03at% 19.3 7.9 69.4 1.4 0.21 1.27 <0.002 0.16 0.266 0.06 0.050

3.2 Proton Irradiation

Samples for post-irradiation annealing studies were irradiated with 3.2 MeV protons to

doses 1.0 or 2.5 dpa at a dose rate of approximately 4 ¥ 10-6 dpa/s, resulting in a nearly uniform

damage rate throughout the first 35 mm of the proton range (40 mm), Fig. 3. The sample

temperature during irradiation was maintained at 360° ± 10°C. A displacement energy of 25 eV


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was used for dose calculations throughout this paper. Further details of the sample preparation

and irradiation are given elsewhere. [30] Following irradiation, samples were annealed in a

small tube furnace in a flowing Ar gas stream at temperatures ranging from 400ºC to 650ºC fortimes of 45 or 90 minutes and then quenched in water.

Low temperature irradiations for microstructure analysis were conducted to doses of 0.25

to 1.5 dpa dpa with 3.2 MeV protons at a temperature of no greater than 75°C. A reduction in

dose rate to 5 ¥ 10-7 dpa/s was required to maintain the low sample temperature, which was

monitored with three J-type thermocouples and was kept at less than 75°C throughout the

irradiation. Following low temperature irradiation, samples were annealed in a vacuum furnace

at 350°C for either 8 minutes or 15 minutes.

Figure 3. Displacement rate profile for 3.2 MeV protons in stainless steel as calculated by the MonteCarlo program TRIM 90.

3.3 Radiation Induced Segregation

Microchemical analysis was performed before and after annealing using a scanning

transmission electron microscope with energy-dispersive x-ray analysis (STEM/EDS). The

STEM/EDS analysis was performed in a Philips CM200/FEG at Oak Ridge National Laboratory,which produces a probe approximately 1.2 nm in diameter (full-width, half-maximum) while

operating at 200kV. STEM/EDS measurements were performed on “edge-on” grain boundaries

7x10-6 dpa/s40 hours irrad..


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so as to minimize broadening of the boundary profile. Details of the grain boundary

measurement technique are given in ref. [31].

3.4 Microstructure Analysis

The dislocation microstructures of samples irradiated at 75°C and 360°C, both before and

after annealing were analyzed using JEOL 2010F/FEG and JEOL 4000 instruments at the North

Campus Electron Microscopy Analysis Laboratory at the University of Michigan and the

CM200/FEG at ORNL. For bright-field imaging (BF), a two-beam condition at g=[200] (close

to the <110> zone axis) was used. The dark-field, rel-rod technique was also utilized. Most

dislocation loop images were taken at magnifications of 100-200kx.

3.5 Post-irradiation Hardness Measurement

Hardening for the proton-irradiated alloys was measured using a Vickers hardness indenter(MICROMET II) with a load of 25 g. This load was used to confine the plastic zone ahead of the

indenter tip to a depth within the proton range (~40 mm) to ensure that unirradiated material is not being

sampled. The yield strength of the proton-irradiated heats is a useful parameter for comparison toexisting literature data on changes in yield strength. While yield strength cannot be determined directly

from the proton-irradiated samples, correlations have been developed which allow calculation of

expected yield strength from dislocation microstructure or hardness. The yield strength change

associated with irradiation can be estimated using

Dsy = 3.55 DHv for doses < 2.0 dpa (1)

Dsy = 2.69 DHv for doses > 2.0 dpa

where Dsy is expressed in MPa and DHv is expressed in kg/mm2 . [32]


19

3.6 Stress Corrosion Cracking Experiments

The constant extension rate tensile (CERT) tests were conducted in a multiple-specimen

CERT test system, supplied by Korros Data. The Korros Data system is capable of straining four

samples in parallel, thus providing identical conditions within a given test. Samples were

strained to failure at a rate of 3 x 10-7 s-1. Details regarding the KorrosData system are published

elsewhere [33].

CERT tests were performed in normal water chemistry (NWC) characterized by a water

temperature of 288°C, water conductivity of 0.2 mS/cm and oxygen content of 2 ppm. The

conductivity and oxygen composition were selected to arrive at a value of the corrosion potential

of about +150 mVSHE, representative of BWR cores. [34] The dissolved oxygen concentration

was controlled at 2 ppm by bubbling a 5% O2/Ar mixture through the water reservoir.

Conductivity was controlled via automatic additions of dilute H2SO4 so that the outletconductivity was maintained at 0.2 mS/cm. The electrochemical potential was verified for the

NWC environment. A Cu/CuO reference electrode with a yttria-stabilized zirconia membrane

was used in conjunction with an EG&G Model 173 Potentiostat. A spare CP-304 tensile samplewas used as the working electrode. For the NWC described, the measured potential was +140

mVSHE.

Fractography was performed following each CERT test using a Philips XL30/FEG SEM.

While qualititative fractrography is informative in determining the type of failure and generaltrends, it does not provide any quantititative information about the cracking. Intergranular (IG)

fracture is typically characterized by measurements of the area of IG facets on the fracture

surface and expressed as an area based percentage. However, since proton irradiation only

affects the first 40 mm of the irradiated face, the majority of the fracture surface is unirradiated

material. Therefore, the IG cracking fraction (or percentage) for these samples refers to the

irradiated area (40 mm) only.


20

4.0 RESULTS AND DISCUSSION

The project results will focus on post-irradiation annealing to isolate RIS and on the useof low temperature irradiation to isolate the irradiated microstructure. Section 4.1 covers the

simulation of post-irradiation annealing and section 4.2 covers the experimental results. Section

4.3 provides a discussion of results of simulation and experiment to isolate RIS. Section 4.4

covers the formation of an irradiated microstructure without RIS by low temperature irradiation.

4.1 Post-Irradiation Annealing: Simulation Results

During post-irradiation annealing at low to moderate temperatures (< 700°C), the

removal of composition gradients will be governed by the equilibrium vacancy concentration.

The irradiation-induced composition gradients at grain boundaries will drive the motion of

thermal defects during annealing. The modified inverse-Kirkendall (Perks) model developed by

Allen [27] was used to simulate the behavior of composition gradients during post-irradiation

annealing of 304 SS alloys. The modified inverse-Kirkendall (MIK) model is capable of

handling up to three major alloying elements, and thus, was used to simulate the annealing

behaviors of only Cr, Fe, and Ni. For the HP-304L alloy with the nominal composition listed in

Table 1, the measured segregation profile at 1.0 dpa was used as the initial condition for

annealing simulations. To simulate annealing conditions, the displacement rate was simply set

equal to zero. No other modifications to the model were necessary. Annealing of segregation

profiles was simulated over a wide range of temperatures (350°C to 600°C) and times (up to 107

sec). The annealing of grain boundary Cr depletion is shown in Figure 4 as a function of time for

anneals at 400°C, 500°C, and 600°C. Simulations indicate that at least 106 sec and ~ 1 h at

400°C and 500°C, respectively, are required to cause grain boundary depletion to be reduced to

90% of the as-irradiated condition. However, for annealing at 600°C, 10% of the as-irradiated

degree of depletion is removed in only 30 seconds.


21

Figure 4: Comparison of simulated annealing of fraction of as-irradiated grain boundary Crdepletion and total loop line length as a function of time for anneals at 400°C, 500°C, and600°C. Simulation for CP-304 irradiated to 1.0 dpa at 360°C.

4.1.1 Simulated annealing of dislocation loops and hardness

During post-irradiation annealing, faulted, interstitial loops will absorb thermal vacanciesand shrink in size. The rate of absorption, and hence rate of change in loop size, is also affected

by the line tension and stacking fault energy. In order to simulate the effects of post-irradiation

annealing for a population of interstitial dislocation loops, a model was developed to calculate

0

20

40

60

80

100

1 10 100 1000 104 105 106 107

Annealing Time (sec)

RIS

Loops

Annealing SimulationHP-304L

600oC

1.0 dpa at 360oC

0

20

40

60

80

100

1 10 100 1000 104 105 106 107


RIS

Loops

500oCAnnealing SimulationHP-304L1.0 dpa at 360

oC

0

20

40

60

80

100

1 10 100 1000 104 105 106 107


RIS

Loops

400oCAnnealing SimulationHP-304L1.0 dpa at 360

oC


22

the changes in loop radius and density as a function of time at any given temperature. For a

population of defects widely spaced in comparison to their size, the vacancy diffusion field

around each loop is assumed to be spherically symmetrical. The diffusion equation for motion ofvacancies to a loop is given as:

(2)

,

where r is the radial distance from the dislocation loop, D is the vacancy diffusion coefficient,and C is the vacancy concentration. This equation can be solved by integration using the

boundary conditions; C = Cd at r = rd and C = Ceq at r = R, where rd is the loop radius, Cd is the

vacancy concentration at the loop, R is a characteristic distance, and Ceq is the thermal

equilibrium concentration of vacancies. The flux of vacancies at the loop can be written as:

(3)

.

Following the methodology used by Burton [35], the rate of change of loop radius is given by:

(4)

where U (ªGb2/2) is the dislocation line energy, G is the stacking fault energy, k is the

Boltzmann constant, T is the absolute temperature, G is the shear modulus, and b is the

magnitude of the Burgers vector. Note that the form of Eq. (4) indicates that larger loops are

removed more slowly than smaller loops. This is reasonable, as larger loops will be more stable(i.e. less excess free energy per interstitial) and more vacancies are required to annihilate the

interstitials within a larger loop.

In order to simulate the annealing behavior of a population of dislocation loops, Eq. (4)

was applied to every size group within a population. For a given time step, the amount of change

in loop radius and the radius at the end of the time step were calculated for each group in thepopulation. The model then iterates over time and recalculates the rate of radius change for each

group. When the radius of any group shrinks below 1 nm, the group is removed from the

= -2b2CeqD{1 - exp(-(U/br + G/b)b3/kT)}dr

dt

dn( )r=rd= -4pr2

ddt

dC

dr= 4prdD (Cd-Ceq)

d(r 2 D ) = 0

dr

dC

dr


23

population, thereby reducing the dislocation loop density. While loop size distributions are

informative, quantities such as total loop line length are more meaningful as they give a more

complete description of the dislocation sink strength in the irradiated microstructure. Total loopline length, Sl, can be calculated as:

(5)

In a similar fashion, ÷Nd (which is proportional to the change in yield stress according to the

dispersed barrier hardening model [36]) can also be calculated.

The dislocation loop population of the HP-304L alloy irradiated with protons to 1.0 dpaat 360°C was used as the starting condition for a series of simulated anneals over a wide range of

times and temperatures. The annealing of this dislocation loop population is shown as a function

of time in Figure 2 for anneals at 400°C, 500°C, and 600°C. As the temperature increases from

400°C to 600°C, the time to remove 10% of the as-irradiated loop population drops from 105 sec

to ~10 sec.

The removal of RIS and dislocation loops during post-irradiation annealing is compared

directly in Figure 4. Clearly, the simulations indicate that the dislocation microstructure is

removed preferentially. In order to determine the origin of the difference in the simulatedremoval rates, both thermodynamic and kinetic processes are considered. Specifically, the

apparent activation energy is determined for the removal of both RIS and dislocation loops.

Both the density of vacancies required and potential competition for thermal vacancies between

RIS and dislocation loops are also considered.

4.1.2 Thermodynamic considerations

The removal of both loops and RIS reduces the excess free energy of the system. If

dislocation loop formation during irradiation increases the free energy of the system more than

that due to the presence of segregation profiles, the driving force for dislocation loop removal

will be greater than that for RIS. Since the annealing of both RIS and dislocation loops is

dependent upon the formation and diffusion of vacancies, it may be possible to gain some insight

into the energetics of the removal processes. Removal of dislocation loops and RIS during post-

Slii

. N

S = 2 p

r

All groups,i


24

irradiation annealing is analogous to the removal of cold work during the recovery stage of

annealing which follows a simple activation or Arrhenius type law:

. (6)

This type of relationship was used to determine the activation energy for removing a dislocation

loop population and segregation profiles. The time for removal of 90% of either, t90, was

determined from simulations as a function of post-irradiation annealing temperature. Simulation

of microstructure annealing was accomplished using Eq. (4) and simulation of RIS was

performed using the MIK model. The natural logarithm of 1/t90 is plotted as a function of the

inverse-temperature for the annealing of both RIS and dislocation loops in Figure 5 and the slope

of the least-squares fit gives the apparent activation energy. From Figure 5, the apparent

activation energies for removal of RIS and dislocation loops are 3.1 and 2.5 eV, respectively.

The removal of both dislocation loops and RIS profiles is dependent on diffusion

processes. The apparent activation energy determined from Eq. (6) and Figure 5 for removal of

RIS profiles is 3.1 eV. No combination of input diffusion parameters to the MIK model matches

this value exactly, but the vacancy formation and migration energies for Ni are the closest (1.79

and 1.04 eV, respectively). [27] The total energy required for motion via Ni atoms is 2.83 eV.

While lower than the calculated value, Ni-vacancy diffusion is a sensible mechanism since both

Cr depletion and Ni enrichment profiles must be removed from the grain boundary and the

diffusion controlled process is limited by the rate of diffusion of the slowest participant (in this

case Ni).

The annealing of dislocation loops is also dependent upon the diffusion of vacancies

through the matrix. However, the stacking fault energy and line tension of the dislocation loop

also influence the annealing process by reducing the energy barrier required for removal.

Accounting for the stacking fault energy and line tension results in an apparent activation energy

for loops which is 0.6 to 0.15 eV lower than that for RIS for loops ranging in size from 1 to 20

nm, respectively. The apparent activation energy for loops determined from Figure 5 is 2.5 eV

or 0.6 eV lower than that for removal of RIS.

1

t= A e (-Q/kT)


25

Figure 5: Comparison of activation energies for removal of dislocation loops and RIS duringsimulated post-irradiation anneal of HP-304L irradiated to 1.0 dpa at 360°C.

4.1.3 Kinetic considerations

Although, thermodynamics are important in the annealing processes, kinetic factors may

also play a role. The pre-exponential term in Eq. (6) is also of interest and was also determined

to be 2.2 x 1012 s-1 for loops and 1.1 x 1012 s-1 for RIS. The difference in pre-exponential terms

also gives insight into the importance of the density of defects required for annealing to be

discussed in the next section. Consider the case where annealing is performed at extremely hightemperatures (where 1/kT approaches 0). In this regime, an infinite, inexhaustible supply of

vacancies is available for annealing of both RIS and loops. Yet, a difference in removal rates

still exists, implying kinetics are an important consideration at low temperatures.

The most straightforward comparison between loop and RIS annealing kinetics is the

number of vacancies required to remove a population of dislocation loops versus that for removal

of the segregation profiles. Since the nature of the two irradiation-induced features is different,

-20

-15

-10

-5

0

13 14 15 16 17 18 19

1/kT (eV-1)

Removal of RIS

Removal of Loops


Q = 3.1 eV

Q = 2.5 eV

600 550 500 450 400 350

Annealing Temperature ( oC)

1.0 dpa at 360oC


26

comparisons of the number of vacancies to remove them must be made on a common scale, such

as the number of defects per unit volume.

A 10.5 nm interstitial dislocation loop lying on the (111) plane is a conglomeration of

~3200 interstitials. Therefore, to remove this loop entirely, a net of 3200 vacancies must be

absorbed at each loop. Given a dislocation density of 4.9 x 1015 loops/cm3, a total of 1.6 x 1019

vacancies/cm3 are required to completely remove the dislocation loop population.

The number of vacancies required to remove the segregation profiles per cm3 can also be

calculated. The MIK code was used to track the number of vacancies passing a marker plane set

at 10 nm (maximum width of RIS profile after 1.0 dpa at 360°C). The number of vacancies was

then integrated over time until the segregation profile was completely removed. For the profiles

typical of 1.0 dpa irradiation, a minimum of 3.8 x 107 vacancies per 100 nm2 of grain surface

area are required to remove the segregation profiles. For 11.5 mm grains, there are

approximately 5200 cm2 of grain boundary area per cm3. Thus, a total of 2.0 x 1023

vacancies/cm3 are required to completely remove the segregation profiles. Comparison of the

number of vacancies required for removal indicates that a factor of 2 x 104 more vacancies per

cm3 are required to remove the segregation profiles than the dislocation densities. Clearly, since

fewer vacancies are required to remove dislocation loops than RIS profiles, equal defect fluxes to

each sink will result in more rapid removal of dislocation loops.

Since grain boundaries and dislocation loops are both defect sinks, they may also

compete for the same thermal vacancies during annealing. Specifically, if a vacancy traveling to

a grain boundary passes other sinks, the probability that it reaches the grain boundary is greatly

reduced. Grain boundaries act as a planar sink for defects created throughout the grain, while the

dislocation loop population is a series of sinks spread throughout the grain. Note, however, that

only those vacancies relatively close to the grain boundary will be influenced by the solute

composition gradient and move towards the boundary to participate in annealing. The width of

the vacancy concentration profile during post-irradiation annealing was investigated using the

MIK code. During annealing at 500°C, the vacancy concentration returns to the equilibrium

concentration ~220 nm away from the grain boundary. Therefore, within ~200 nm of the grain

boundary, the boundary and loops in this region compete for vacancies, while in regions more

than 200 nm away, loops are the only sink for vacancies.


27

For the CP-304 SS at 1.0 dpa, the mean loop size is 4.9 nm and the density was 17.0 x

1021 m-3. For a homogeneously distributed population, the distance between loops is, on average,

~110 nm. In this situation, a diffusing vacancy may encounter more than one dislocation loop

over its diffusion path towards the grain boundary, increasing the likelihood that it will be

absorbed at a loop rather than a grain boundary. Another MIK simulation was performed to

evaluate the influence of the dislocation loop density on annealing of RIS profiles. Simulations

of annealing at 500°C were performed with both the as-irradiated dislocation density and with no

loops. The results for the two simulations are compared in Figure 6. Clearly, reducing the

dislocation density increases the rate of annealing of RIS, which confirms that loss of vacancies

to competing sinks such as dislocation loops reduces the annealing rate of RIS.

Figure 6: Simulation of annealing of Cr segregation as a function of time during post-irradiationannealing at 500°C. Simulation with as-irradiated dislocation density and no dislocation density.

0

20

40

60

80

100

1 10 100 1000 104 105 106 107



Annealing at 500oC

As-Irradiated Density (2 x 1014

m-2

)No Dislocation Loops

1.0 dpa at 360oC


28

In summary, the preferential removal of dislocation loops can be explained by

considering both thermodynamic and kinetic factors. Calculation of the apparent activation

energy, the density of defects required for annealing, and potential competition between loopsand RIS for thermal vacancies all support the preferential removal of dislocation loops during

post-irradiation annealing.

4.2 Post-Irradiation Annealing: Experimental Results

4.2.1 RISDuring post-irradiation annealing, the degree of Cr depletion did not change significantly

from the irradiated condition until the most extreme annealing conditions for either the HP-304L

at 1.0 dpa or the CP-304 at 1.0 and 2.5 dpa. Similarly, the amount of Ni and Fe segregation did

not change appreciably. However, for the CP-304 alloy the enrichment of minor elements suchas Si and P was removed rapidly with annealing.

HP-304L

The results of composition measurements on samples of the HP-304L irradiated to 1.0dpa in the as-irradiated condition and the post-irradiation annealed condition are listed in Table

2. (in at%). The number of measurements for each condition is given in the left-most column.

Typical Cr segregation profiles for each annealing condition are shown in Figure 7 a.). All plots

in Figure 7 a.) also contain the as-irradiated profile (shown as open symbols) for direct

comparison. Grain boundary Cr depletion remained virtually unchanged with annealingtreatments below treatments of 600°C/45 min. Annealing at 600°C/90 min. removed, on average,

only 17% of the as-irradiated depletion (17.4 at% versus 16.4 at% in the as-irradiated case) andresulted in little change to the shape of the segregation profile while annealing at 650°C/45 min.

removed 66% of the as-irradiated Cr depletion.

Similarly, the amount of grain boundary Ni enrichment also remained virtually unchanged

with annealing treatments up to 650°C/45 min., which removed virtually all as-irradiated Ni

enrichment. The measured grain boundary Fe enrichment remained at or above the amount

measured in the as-irradiated condition (2.5 at% enrichment) with annealing treatments up to


29

annealing treatments of 600°C/45 min. Further annealing removes additional Fe enrichment,

with the grain boundary Fe content dropping to 70.4 at% (1.4 at% enrichment after annealing at

650°C/45 min.).

Table 2: Summary of grain boundary composition measurements on post-irradiation annealed HP-304L (irradiatedto1.0 at 360°C). All results are listed in at%.

Irradiation Condition Fe Cr Ni Mn

HP-304L matrix/bulk comp. 69.0 20.9 9.0 1.1

As-irradiated-1.0 dpa

GB avg. (25 meas.) 71.5 16.4 11.4 0.72

Std. Dev. Of mean (at%) 0.2 0.3 0.2 0.03

500°C/45 min.

GB avg. (28 meas.) 72.1 16.1 11.1 0.75

Std. Dev. Of mean (at%) 0.2 0.2 0.2 0.02

500°C/300 min.

GB avg. (8 meas.) 72.5 16.1 10.8 0.71

Std. Dev. Of mean (at%) 0.5 0.3 0.3 0.04

550°C/45 min.

GB avg. (9 meas.) 71.7 16.6 10.9 0.79

Std. Dev. Of mean (at%) 0.5 0.4 0.3 0.06

600°C/45 min.

GB avg. (18 meas.) 72.1 15.4 11.8 0.75

Std. Dev. Of mean (at%) 0.2 0.3 0.4 0.04

600°C/90 min.

GB avg. (26 meas.) 71.0 17.2 10.8 0.96

Std. Dev. Of mean (at%) 0.2 0.1 0.2 0.05

650°C/45 min.GB avg. 13 meas.) 70.4 19.4 9.2 0.96

Std. Dev. Of mean (at%) 0.5 0.6 0.4 0.05


30

Figure 7. Cr segregation profiles following irradiation with 3.2 MeV protons at 360°C and post-irradiation annealing: HP-30L to 1.0 dpa and CP-304 to 1.0 dpa and 2.5 dpa. The as-irradiatedprofile is shown by the open symbol. The 0 nm position is the grain boundary for all profiles.

CP-304

The composition measurements for the CP-304 samples irradiated to 1.0 and 2.5 dpa in

both the as-irradiated and post-irradiation annealed conditions are summarized in Table 3. (in

at%) and shown in Figures 7 b.) and 7 c.) for the 1.0 and 2.5 dpa samples, respectively. At 1.0

12

14

16

18

20

22

-12 -8 -4 0 4 8 12

Distance from GB (nm)

As-Irrad.

650oC/45 minutesHP-304L 1.0 dpa

12

14

16

18

20

22

-12 -8 -4 0 4 8 12


As-Irrad.


12

14

16

18

20

22

-12 -8 -4 0 4 8 12


As-Irrad.


12

14

16

18

20

22

-12 -8 -4 0 4 8 12


As-Irrad.


12

14

16

18

20

22

-12 -8 -4 0 4 8 12


As-Irrad.

600oC/90 minutesCP-304 1.0 dpa

12

14

16

18

20

22

-12 -8 -4 0 4 8 12


As-Irrad.


12

14

16

18

20

22

-12 -8 -4 0 4 8 12


As-Irrad.


12

14

16

18

20

22

-12 -8 -4 0 4 8 12


As-Irrad.


12

14

16

18

20

22

-12 -8 -4 0 4 8 12Distance from GB (nm)

As-Irrad.


12

14

16

18

20

22

-12-8 -4 0 4 8 12Distance from GB (nm)

As-Irrad.


12

14

16

18

20

22

-12-8 -4 0 4 8 12Distance from GB (nm)

As-Irrad.


12

14

16

18

20

22

-12 -8 -4 0 4 8 12


As-Irrad.



31

dpa, both the degree of Cr depletion and shape of the Cr segregation profiles remained virtually

unchanged for all annealing treatments. In the as-irradiated condition, the Cr segregation profile

has a “w-shape”, that is, the minimum in composition is on either side of the grain boundary witha local maximum in concentration at the grain boundary. Following anneals at 400°C, 450°C,

and 500°C for 45 min., the minimum in measured Cr content near the grain boundary is

unchanged from the as-irradiated level. The Cr content at the grain boundary Cr content

increased slightly with annealing, although, the as-irradiated and all annealed GB Cr contents at

the peak of the “w-shape” were measured at levels above the bulk content. Of more significanceis that the complex, “w-shape” segregation profiles remained unchanged. Even annealing at

600°C/90 min., did not alter the as-irradiated GB Cr content or the segregation profiles.

Similarly, the amount of grain boundary Ni enrichment remained unchanged with all annealingtreatments while the degree of grain boundary Fe depletion was slightly altered with all

annealing treatments as listed in Table 3.

Contrary to Cr, Fe, and Ni, grain boundary Si changes significantly during post-

irradiation annealing as shown in Table 3. The grain boundary Si enrichment was reduced

steadily with annealing, with only 0.2 at% enrichment remaining after annealing at 600°C/90

min. Grain boundary P also changes drastically during post-irradiation annealing. Following

annealing at 400°C/45 min., the grain boundary P content has dropped to nearly the bulk level

(from 1.47 at% in the as-irradiated condition) where it remains for all other annealing treatments.

For the samples irradiated to 2.5 dpa, the grain boundary Cr profiles remained virtually

unchanged with all annealing treatments. Contrary to the 1.0 dpa case the higher dose produces

a more traditional “v-shape” profile and a higher degree of Cr depletion. Annealing at 600°C for

90 min. did not alter the as-irradiated GB Cr content or the segregation profiles for the CP-304

irradiated to 2.5 dpa as illustrated in Figure 7c.). The grain boundary Ni and Fe content

remained relatively unchanged from the as-irradiation condition, similar to the 1.0 dpa samples.

Similar to the 1.0 dpa case, grain boundary Si changed significantly during post-

irradiation annealing with the grain boundary Si content dropping significantly with increasing

annealing time or temperature with only 0.3 at% enrichment remaining after annealing at600°C/90 min. Grain boundary P also changed drastically during post-irradiation annealing.

Following annealing at 400°C for 45 min., the grain boundary P content dropped to 0.46 at%


32

from 1.47 at% in the as-irradiated condition. All other treatments resulted in grain boundary P

contents near the bulk level of 0.055 at%.

Table 3: Summary of grain boundary composition measurements* on post-irradiation annealed CP-304(irradiated to1.0 and 2.5 dpa at 360°C). All results are listed in at%.

Irradiation Condition Fe* Cr* Ni P Mo* Mn Si

CP-304 matrix/bulkcomp.

69.4 19.3 7.9 0.055 0.21 1.40 1.27

As-irradiated 1.0 dpa

GB avg. (50 meas.) 66.7/68.2 19.6/18.5 8.6 1.47 0.41/0.22 1.09 2.14

Std. Dev. Of mean(at%)

0.2/0.2 0.1/0.1 0.1 0.01 0.02/0.02 0.05 0.02

400°C/45 min.

GB avg. (10 meas.) 67.9/68.9 19.8/18.6 8.7 0.08 0.41/0.35 1.27 1.85


0.5/0.5 0.3/0.3 0.2 0.01 0.1 0.11 0.05

450°C/45 min.

GB avg. (20 meas.) 67.9/69.1 20.2/18.6 8.5 0.07 0.27/0.27 1.18 1.79


0.3/0.3 0.3/0.3 0.2 0.01 0.01 0.11 0.05

500°C/45 min.

GB avg. (21 meas.) 67.4/69.5 20.9/18.9 8.5 0.07 0.27/0.23 1.17 1.67


0.3/0.3 0.2/0.2 0.1 0.01 0.02/0.02 0.02 0.05

600°C/90 min.

GB avg. (10 meas.) 67.9/69.2 20.3/18.8 8.9 0.06 0.22/0.22 1.13 1.47


0.2/0.2 0.1/0.1 0.2 0.01 0.01/0.01 0.07 0.06

As-irradiated 2.5 dpa

GB avg. (20 meas.) 66.8 18.4 9.7 1.47 0.35 0.85 2.42


0.2 0.2 0.2 0.03 0.03 0.05 0.05

400°C/45 min.

GB avg. (10 meas.) 68.0 18.3 9.9 0.46 0.33 1.00 2.02


0.3 0.1 0.2 0.01 0.1 0.04 0.05

450°C/45 min.

GB avg. (13 meas.) 68.2 18.4 9.7 0.08 0.31 0.82 1.95


0.5 0.1 0.3 0.01 0.05 0.05 0.07

500°C/45 min.

GB avg. (21 meas.) 68.8 18.4 9.8 0.07 0.25 0.83 1.77


0.3 0.1 0.2 0.01 0.01 0.03 0.05

600°C/90 min.

GB avg. (14 meas.) 68.8 18.3 9.7 0.07 0.22 0.91 1.57


0.4 0.1 0.3 0.02 0.02 0.06 0.08*Average at GB/Max (Min) segregation measured adjacent to boundary for “W-shaped” profiles.


33

The removal of Cr segregation for the HP-304L and CP-304 irradiated to 1.0 and 2.5 dpa

is plotted in Figure 8 as a function of Fe-diffusion distance. Because both time and temperature

were varied, the diffusion distance of Fe was chosen as a single variable which combines bothtime and temperature effects. The diffusion distance is equal to (DFet)1/2 where t is annealing time

and DFe is diffusivity for Fe at the annealing temperature. Also plotted in Figure 8 are the results

from experimental studies of Jacobs et al. [37], Katsura et al. [38], and Bruemmer et al. [39] for

CP-304 and 316 irradiated to ~1.0 dpa. For alloys and doses where a “w-shape” profile was

observed, the percent of the minimum measured value remaining is plotted. Note that forconditions where the amount of Cr segregation increased with annealing, the percentage

remaining is plotted as 100%. This is most significant in the data of Bruemmer where a CP-316

alloy with an extremely sharp “w-shape” profile was measured. The Cr content measured in the

as-irradiated case was above the bulk level and continued to increase above the bulk level duringpost-irradiation annealing. Overall, the data from this study and other studies are in excellent

agreement. Only two annealing conditions (HP-304L at 650°C/45 min. and Katsura’s CP-316 at

650°C/1 h) resulted in the Cr segregation being less than 80% of the as-irradiated value.

Figure 8. Annealing of Cr segregation as a function of Fe-diffusion distance. The % of as-irradiated minimum measured Cr is plotted for all conditions. Data points for CP-304 at 1.0 dpaand 2.5 dpa have been shifted left and right, respectively, for clarity.

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002 0.0025 0.003

Iron Diffusion Distance, (DFe

t)1/2, (cm)

Data Trend

HP-304L (1.0 dpa)

Jacobs 304 (~2.6 dpa)

Bruemmer 316 (~1.7 dpa)

CP-304 (1.0 dpa)

CP-304 (2.5 dpa)

Nishimura 316 (~11 dpa)


34

In Figure 9, the experimental data are compared to the predictions of the MIK model.

The model predictions are plotted as a function of the measured result from both this study and

available literature data. When plotted in this fashion, model and measured results can be

compared for all available data. A line with a slope of unity was applied to the data set (shown

in Figure 9), resulting in a correlation coefficient of 0.85. A line with slope = 1 represents a one-

to-one correlation between modeled and measured results. However, the most significant

differences between the simulation and measured data occur where the measurements show no

change in the as-irradiated Cr depletion. For these data points the model clearly overpredicts the

rate of annealing. The majority of these data points are for conditions containing “w-shape”

profiles (CP-304 at 1.0 dpa and CP-316 data from Bruemmer), which the MIK model is not

capable of simulating accurately [40]. Excluding these points from the statistical analysis

improves the correlation coefficient to 0.93.

Figure 9. Comparison of measured annealing of Cr segregation with simulated annealing usingthe MIK model.

0

20

40

60

80

100

0 20 40 60 80 100

Measured % of as-irradiated Cr depletionremaining after heat treatment

HP-304L (1.0 dpa)


Bruemmer 316 (~1.7 dpa)

CP-304 (1.0 dpa)

CP-304 (2.5 dpa)

Katsura 316 (~11 dpa)

For all data: R2 = 0.85

Without "w-shape profiles": R2 = 0.93


35

Grain boundary segregation of Si and P is removed much faster than Cr or Ni for the

same annealing treatments. For both the 1.0 and 2.5 dpa CP-304 samples, the as-irradiated

enrichment of Si, and P are all significantly affected by annealing at 400°C for 45 min., contrary

to the observed annealing of Cr and Ni. The more rapid removal of Si segregation during post-

irradiation annealing can be explained by considering the tracer impurity diffusion coefficients

[41]. Silicon diffuses considerably faster than Cr or Ni. Coupled with a considerable Si

enrichment in the 1.0 and 2.5 dpa condition (1.7 and 1.9 times bulk content, respectively), Si

enrichment should be removed rapidly during post-irradiation annealing. Phosphorous also

diffuses via vacancies in Fe-Cr-Ni alloys [42], although the diffusion coefficients are less well

known than those for the major and other minor alloying elements. Nonetheless, in the as-

irradiated 1.0 and 2.5 dpa condition, P is enriched a factor of 27 times over the bulk content (1.47

at% for both 1.0 and 2.5 dpa versus a bulk content of 0.055 at%). The extreme concentration

gradient in the as-irradiated condition can explain the very rapid removal of P enrichment during

post-irradiation annealing.

4.2.2 Dislocation loops

During post-irradiation annealing, dramatic changes were measured in the dislocation

loop population. For the HP-304L at 1.0 dpa and CP-304 to 1.0 and 2.5 dpa, the dislocation loop

density decreased steadily with increasing annealing time or temperature while the meandislocation size remained relatively unchanged. The mean loop diameters and loop densities

were determined for each annealing condition and are summarized in Table 4 for both alloys.

Bright field images of the dislocation population before and after annealing are shown in Figure

10.

HP-304L

For the HP-304L, in the as-irradiated condition, the mean loop diameter was 11.0 nm.

Annealing at 500°C for 45 min. resulted in a slightly smaller mean diameter (10.8 nm) while

anneals at 600°C for 45 min., and 600°C for 90 min. both resulted in a slightly larger loop

diameter (11.7 and 12.1 nm, respectively). Bright field images of the dislocation loops in both


36

the as-irradiated and annealed conditions are shown in Figure 10a. Some dislocation loops that

have unfaulted and grown (not discriminated by the bright field imaging) may have contributed

to this measured growth. Also, the simulations predicted the preferential removal of small loopsfrom the population that would also result in an apparent growth in mean loop diameter.

Table 4. Summary of dislocation loop analysis on post-irradiation annealed HP-304L and CP-304 (irradiated to1.0 and 2.5 dpa at 360°C).

Annealingtemp. (°C)

Annealingtime (min.)

Mean loopdiameter

(nm)

Loopdensity

x 1021 (# /m3)

Total loopline length

x 1014

(m/m3)

% As-irradiated line

lengthremaining

HP-304L 1.0 dpa

As-irradiated 11.0 5.6 1.94 100.0

500 45 10.8 4.2 1.43 73.6

600 45 11.7 1.9 0.69 36.1

600 90 12.1 0.16 0.06 3.1

650 45 9.2 0.14 0.04 2.1

CP-304 1.0 dpa

As-irradiated 4.9 17.0 2.61 100.0

400 45 4.8 18.0 2.71 100.0

450 45 5.6 16.5 2.90 100.0

500 45 5.8 12.0 2.19 83.8

600 90 6.7 0.29 0.06 2.4

CP-304 2.5 dpa

As-irradiated 5.2 40.0 6.53 100.0

400 45 5.4 38.7 6.56 100.0

450 45 5.4 37.1 6.29 96.4

500 45 5.8 18.7 3.41 52.2

600 90 6.1 0.96 0.18 2.8


37

Alternatively, a portion of the dislocation loop population may, indeed, be composed of

vacancy-type dislocation loops, that will grow with annealing. Computer simulation of the post-

irradiation annealing of both vacancy and interstitial type dislocation loops by Simonen et al.

[43] suggests that both the presence of vacancy-type loops in a loop population and their growthduring annealing are feasible. Only annealing at 650°C/45 min. resulted in a loop size

significantly smaller (9.2 nm) than the as-irradiated value. Contrary to the mean loop diameter,

the dislocation loop density dropped significantly during annealing. In the as-irradiatedcondition, a loop density of 5.6 x1021m-3 was measured, which was steadily reduced with

annealing. Following the most severe annealing condition (650°C for 45 min) a density of only

0.14 x1021m-3 remained.

Figure 10. Bright field images of dislocation loop populations in HP-304L and CP-304irradiated with 3.2 MeV protons at 360°C to 1.0 dpa and 2.5 dpa and post-irradiation annealed.

rd = 0.96 x 1021 m-3

d = 6.1 nm 50 nm

rd = 40.0 x 1021 m-3

d = 5.2 nm 50 nm rd = 37.1 x 1021 m-3

d = 5.4 nm 50 nm rd = 18.7 x 1021 m-3

d = 5.8 nm 50 nm rd = 38.7 x 1021 m-3

d = 5.4 nm 50 nm

rd = 0.29 x 1021 m-3

d = 6.7 nm 50 nm

rd = 17.0 x 1021 m-3

d = 4.9 nm 50 nm rd = 16.5 x 1021 m-3

d = 5.6 nm 50 nm rd = 12.0 x 1021 m-3

d = 5.8 nm 50 nm rd = 18.0 x 1021 m-3

d = 4.8 nm 50 nm

600°C/90 min.500°C/45 min.450°C/45 min.400°C/45 min.As-Irrad.

rd = 0.14 x 1021 m-3

d = 9.2 nm 25 nm

rd = 5.6 x 1021 m-3

d = 11.0 nm 50 nm rd = 1.9 x 1021 m-3

d = 11.7 nm 50 nm rd = 0.16 x 1021 m-3

d = 12.1 nm 25 nm rd = 4.2 x 1021 m-3

d = 10.8 nm 50 nm


CP-

304

2.5

dpa

CP-

304

1.0

dpa

HP -

304L

1.0

dpa

rd = 0.96 x 1021 m-3

d = 6.1 nm 50 nm

rd = 40.0 x 1021 m-3

d = 5.2 nm 50 nm rd = 37.1 x 1021 m-3

d = 5.4 nm 50 nm rd = 18.7 x 1021 m-3

d = 5.8 nm 50 nm rd = 38.7 x 1021 m-3

d = 5.4 nm 50 nm

rd = 0.29 x 1021 m-3

d = 6.7 nm 50 nm

rd = 17.0 x 1021 m-3

d = 4.9 nm 50 nm rd = 16.5 x 1021 m-3

d = 5.6 nm 50 nm rd = 12.0 x 1021 m-3

d = 5.8 nm 50 nm rd = 18.0 x 1021 m-3

d = 4.8 nm 50 nm


rd = 0.14 x 1021 m-3

d = 9.2 nm 25 nm

rd = 5.6 x 1021 m-3

d = 11.0 nm 50 nm rd = 1.9 x 1021 m-3

d = 11.7 nm 50 nm rd = 0.16 x 1021 m-3

d = 12.1 nm 25 nm rd = 4.2 x 1021 m-3

d = 10.8 nm 50 nm


rd = 0.96 x 1021 m-3

d = 6.1 nm 50 nm

rd = 40.0 x 1021 m-3

d = 5.2 nm 50 nm rd = 37.1 x 1021 m-3

d = 5.4 nm 50 nm rd = 18.7 x 1021 m-3

d = 5.8 nm 50 nm rd = 38.7 x 1021 m-3

d = 5.4 nm 50 nm

rd = 0.29 x 1021 m-3

d = 6.7 nm 50 nm

rd = 17.0 x 1021 m-3

d = 4.9 nm 50 nm rd = 16.5 x 1021 m-3

d = 5.6 nm 50 nm rd = 12.0 x 1021 m-3

d = 5.8 nm 50 nm rd = 18.0 x 1021 m-3

d = 4.8 nm 50 nm


rd = 0.14 x 1021 m-3

d = 9.2 nm 25 nm

rd = 5.6 x 1021 m-3

d = 11.0 nm 50 nm rd = 1.9 x 1021 m-3

d = 11.7 nm 50 nm rd = 0.16 x 1021 m-3

d = 12.1 nm 25 nm rd = 4.2 x 1021 m-3

d = 10.8 nm 50 nm


CP-

304

2.5

dpa

CP-

304

1.0

dpa

HP -

304L

1.0

dpa


38

CP-304

For the CP-304 irradiated to 1.0 dpa and annealed, the very dense dislocation populationin the as-irradiated condition is steadily reduced with annealing, as indicated in the bright field

images shown in Figure 10b. In the as-irradiated condition, the mean loop diameter was 4.9 nm,

but at a density of 17.0 x1021m-3. During annealing, the mean loop diameter increased slightly

(up to 6.7 nm after 600°C/90 min). The dislocation loop density steadily decreased during

increased annealing time or temperature. Following annealing at 600°C for 90 min., very few

loops were observed.

Bright field images of the dislocation population following proton irradiation of the CP-

304 to 2.5 dpa and subsequent annealing are also shown in Figure 10. As in the samples

irradiated to 1.0 dpa, the very dense dislocation population in the as-irradiated condition is

steadily removed with annealing. The mean loop size grew from the as-irradiated diameter of5.2 nm to 6.1 nm following annealing a 600°C/90 min anneal. Loop density dropped from the

as-irradiated density of 40.0 x1021m-3 to 0.96 x1021m-3 after the same 600°C/90 min anneal.

The behavior of the dislocation microstructure during post-irradiation annealing of the

CP-304 alloy in this study is summarized in Figure 11. Also plotted are the results for neutron-

irradiated 304SS by Jacobs et al. [37]. The total dislocation line length associated with the

dislocation loop population is plotted as a function of Fe-diffusion distance. The loop line length

from the proton-irradiated samples is removed steadily with increasing annealing time or

temperature. Jacob’s data, however, are somewhat contradictory. The two data points from the

neutron-irradiated samples indicate that the line length actually increases during annealing.

Jacobs measured a large increase in loop diameter (6.7 to 10.7 nm) for the anneal at a diffusion

distance of 0.0017 cm (475°C for 24 h) and attributed this discrepancy to the unfaulting of

dislocation loops which artificially increased the mean loop diameter.

The experimental data of both this study and that of Jacobs are compared to the

dislocation loop-annealing model in Figure 12. The simulated results are plotted as a function of

measured results for each experimental data point in this study and available data from other

studies. A line with a slope of one is also plotted in Figure 10. This line fits the data set with a

correlation coefficient of 0.92. As with the comparison of RIS annealing shown in Figure 9,

there is no systematic difference between measured and model results.


39

Figure 11. Annealing of dislocation microstructure as a function of Fe-diffusion distance. Thefraction of the as-irradiated loop line length associated with the dislocation population is plotted.

Figure 12. Comparison of simulated and measured annealing of dislocation loop line length.

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002 0.0025 0.003


t)1/2, (cm)

HP-304L (1.0 dpa)


CP-304 (1.0 dpa)

CP-304 (2.5 dpa)

RISData Trend

0

20

40

60

80

100

0 20 40 60 80 100

Measured % of as-irradiated loop line lengthremaining after heat treatment

HP-304L (1.0 dpa)


CP-304 (1.0 dpa)

CP-304 (2.5 dpa)

R2

= 0.92


40

However, comparison of the measured loop sizes with the simulated results reveals a

discrepancy. The loop annealing model predicts that the mean loop size will shrink steadily

during annealing while experimental results show that the loop size remains constant or increases

slightly during annealing up to anneals at 600°C for 90 min., beyond which the loop radius

decreases. Thus, in order for the simulated change in yield stress to match the experimental

values, the model must underpredict the change in loop density to compensate for the

overprediction of change in loop radius. This discrepancy may be explained by the unfaulting of

dislocation loops during annealing, as noted by Jacobs. Dislocation loops that unfault are free to

grow or glide and will behave differently during post-irradiation annealing. According to

Olander [44], unfaulting is very slow at temperatures below ~550°C, but dislocation loops may

spontaneously unfault at temperatures above 600°C. Since the bright field imaging technique

used in this study images both faulted and unfaulted dislocation loops, the reported diameter and

density may have been determined from a population containing some unfaulted dislocation

loops, similar to the experience of Jacobs. The dark-field rel-rod technique images only faulted

dislocation loops and could resolve this issue. However, the rel-rod technique is extremely hard

to use to image very low dislocation densities, which is why the bright field technique was used

exclusively in this study.

4.2.3 Hardness

As with dislocation loop density, the measured hardeness decreases steadily with

increasing annealing time or temperature for both the HP-304L and the CP-304 (at both 1.0 and

2.5 dpa). Table 5 lists the results of the hardness measurements for both the as-irradiated and

annealed specimens. For the HP-304L alloy irradiated to 1.0 dpa, a hardness of 229 kg/mm2 (62

kg/mm2 above the unirradiated value) was measured. After annealing at 500°C for 45 min.

hardness was reduced to 212 kg/mm2 (a 28.5% reduction in radiation-induced hardness).

Increased annealing temperature resulted in an increased removal of radiation-induced

hardening. After annealing at 600°C /90 min. or 650°C/45 min. all radiation-induced hardening

was removed.


41

Table 5: Summary of hardness analysis on post-irradiation annealed CP-304 (irradiated to 1.0 and 2.5 dpa at360°C).

From Hardness From Loops

Annealtemp.(°C)

Anneal.time

(min.)

UnirradHv

(kg/mm2)

As-irradHv

(kg/mm2)

As-irrad.DHv

(kg/mm2)

AnnealedHv

(kg/mm2)

AnnealedDHv

(kg/mm2)

Dsy(MPa)

% As-Irrad.Dsy

remain.

Dsy

(MPa)% As-irrad.Dsy

remain.

HP-304L 1.0dpa

As-irradiated 167 229 62 229 62 219.8 100.0 186.2 100.0

500 45 167 223 56 212 45 159.6 72.5 159.8 85.8

500 300 167 234 67 207 40 141.8 59.7 -- --

550 45 167 234 67 217 50 177.3 74.6 -- --

600 45 167 224 57 182 15 53.2 24.2 111.8 60.1

600 90 167 250 83 166 -1 -3.5 -1.2 33.0 17.7

650 45 167 255 88 156 -11 -39.0 -12.5 26.9 14.5

CP-304 1.0 dpa

As-irradiated 220 270.5 50.5 270.5 50.5 179.1 100.0 227.5 100.0

400 45 220 270.5 50.5 273.0 53.0 187.9 105.0 223.4 98.2

450 45 220 270.5 50.5 273.7 53.7 190.4 106.3 215.9 94.9

500 45 220 270.5 50.5 266.4 46.4 164.5 91.9 207.9 91.4

600 90 220 270.5 50.5 228.8 8.8 31.2 17.4 36.8 16.2

650 45 220 270.5 50.5 218.9 -1.1 3.9 -2.2 -- --

CP-304 2.5 dpa

As-irradiated 220 326.0 106.0 326.0 106.0 375.9 100.0 359.5 100.0

400 45 220 326.0 106.0 320.8 100.8 357.4 95.1 360.1 100.0

450 45 220 326.0 106.0 316.1 96.1 340.8 90.6 352.6 98.0

500 45 220 326.0 106.0 290.8 70.8 251.1 66.8 259.9 72.0

600 90 220 326.0 106.0 234.3 13.4 47.5 12.6 60.0 16.7


42

For the CP-304 alloy in the 1.0 dpa as-irradiated condition, a hardness of 271 kg/mm2 (51

kg/mm2 above the unirradiated value of 220 kg/mm2) was measured by Was et al. [29] and

confirmed in this study. After annealing at 400°C/45 min. and 450°C/45 min., the hardness

remained at the as-irradiated level. However, annealing at 600°C/90 min. removed almost all the

as-irradiated hardness increase. Similarly, for the 2.5 dpa samples, annealing at 400°C/45 min.

and at 450°C/45 min. the hardness was not significantly different from the as-irradiated level,

while annealing at 600°C/90 min. removed a more significant portion of the radiation-induced

hardening (only 16.7% remaining).

Figure 13: Annealing of measured hardness as a function of Fe-diffusion distance. The fractionof the as-irradiated change in yield stress calculated from hardness measurements is plotted.

The fraction of radiation-induced yield stress change calculated from hardness

measurements is plotted as a function of Fe-diffusion distance in Figure 13 for both the CP304

irradiated to 1.0 and 2.5 dpa and the HP-304L of this study. The results from studies by Jacobs,

Asano, Bruemmer, and Katsura on 304 SS and Katsura on 316 SS are also plotted. With the

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002 0.0025 0.003


t)1/2, (cm)

RISData Trend

HP-304L (1.0 dpa)


CP-304 (1.0 dpa)CP-304 (2.5 dpa) Bruemmer 316 (~1.7 dpa)

Katsura 316 (~11 dpa)Katsura 304 (~11 dpa)

Asano 304 (~17 dpa)

LoopsData Trend


43

exception of single data points from Bruemmer and Jacobs each, the hardness decreases steadily

with increasing annealing time and temperature for all alloys. Further, the data from this study

are in excellent agreement with that from Jacobs [37], Katsura [138], and Bruemmer [39].

However, more importantly, all the data from hardness measurements follow the same trend

during annealing and are also in excellent agreement with that calculated from microstructure

measurements.

4.2.4 IASCC of Post-Irradiation Annealed Stainless Steel

The strain to failure and extent of intergranular cracking changed dramatically during

post-irradiation annealing for both 1.0 and 2.5 dpa CP-304 samples tested in the water

environment. The strain-to- failure and measured ultimate tensile strength (UTS), as well as the

location and nature of failure for each specimen in CERT tests of the 1.0 and 2.5 dpa samples are

summarized in Table 6. With annealing, the strain-to-failure increases with increasing diffusion

distance for both 1.0 and 2.5 dpa samples. The 1.0 dpa sample annealed at 400°C for 45 min.

failed at 23.3%, close to the strain-to-failure of the as-irradiated specimens, while all other

annealed 1.0 dpa samples failed at ~30% strain, which is more representative of the 0.3 dpa

samples of the same alloy which failed via ductile rupture [29]. Similarly, the strain-to-failure

for the 2.5 dpa sample annealed at 400°C for 45 min. was 26.5% while the sample annealed at

500°C for 45 min. failed at 33.7% strain.

All four 1.0 dpa specimens failed in the unirradiated region or the threads of the sample.

All faces of each specimen and the fractured ends were examined in detail for evidence of IG

cracking. The 1.0 dpa samples annealed at 400°C for 45 min. and 450°C for 45 min. both had

one crack on the irradiated face of the specimen, approximately 250 mm long and IG in nature.

This is approximately the same length of IG cracking observed on the fractured end in the 1.0

dpa as-irradiated specimens strained under the same conditions [29]. For the samples annealed

at 500°C/45 min. and the 600°C/90 min., no cracks were found on any of the sample surfaces.

All four of the post-irradiation annealed 2.5 dpa samples failed in the irradiated region. The

samples annealed at 400°C and 450°C for 45 min. both exhibit a fracture morphology similar to

the 3.0 dpa as-irradiated fracture morphology. In addition to the crack leading to failure, 2 and 6

additional cracks were found on the irradiated surface for the samples annealed at 400°C and


44

450°C, respectively. Contrary to the samples annealed at 400°C and 450°C for 45 min., the

sample annealed at 500°C/45 min. failed entirely via ductile rupture. Finally, the 2.5 dpa sample

annealed at 600°C/90 min. exhibited extensive IG cracking on all four faces of the sample, not

just the irradiated face, similar to that observed in sensitized specimens. Cookson [45] tested a

sensitized HP-304L sample (650°C for 24 h) under identical water conditions and observed

almost 100% IG failure at only 11.5% strain (compared to 30% strain-to-failure for unirradiated

specimens), suggesting that the 2.5 dpa sample may have actually been annealed at temperature

higher than 600°C and sensitized (potentially due to a thermocouple failure).

Table 6. Summary of CERT test results performed on post-irradiation annealed CP-304 samples(1.0 and 2.5 dpa).

Annealingtemp. (°C)

Annealingtime

(min.)

Strain atfailure

(%)

UTS(ksi)

Location offailure

Numberof cracks

inirradiated

region

Total cracklength onirradiatedsurface(mm)#

CP-304 1.0 dpa

-- --

400 45 23.3* 62.1** Unirrad.region

1 270

450 45 30.0* NA Unirrad.region

1 234

500 45 30.6 66.6 Threads 0 0

600 90 30.6 75.8 Near shoulder 0 0

CP-304 2.5 dpa

400 45 26.5 60.3 Irrad. region 3 1540

450 45 22.7 59.1 Irrad. region 7 3540

500 45 33.7 64.0 Irrad. region 0 0

600 90 10.8 40.5 Irrad. region 0 0

#Sum of length of all regions characterized as IG or TG, including crack that led to failure.*ESTIMATED DUE TO LOAD CELL FAILURE PRIOR TO SAMPLE FAILURE. ± 0.5%**Maximum measured before load cell failed.


45

The percentage of as-irradiated crack length remaining after post-irradiation annealing is

plotted in Figure 14. Even the lowest temperature anneal reduced the crack length relative to the

as-irradiated level. Annealing at 500°C for 45 min. and above removed all evidence of IG

cracking in both the 1.0 dpa and 2.5 dpa samples. The rapid and distinct change in cracking

mode from clearly IG after annealing at 450°C to completely ductile after annealing at 500°C is

also shown in Figure 14. The cracking results of this study are in good agreement with the

cracking response measured by Jacobs [37] and Katsura [38] on neutron-irradiated and annealed

304 and 316 stainless steels, respectively.

Figure 14. Comparison of measured total crack length or %IG remaining as a function and Fe-diffusion distance. Also, the irradiated surface of CP-304-2.5 dpa sample post-irradiationannealed at 450°C and 500°C for 45 min are shown. Samples strained at 3 x 10-7 s-1 in water at288°C, 0.2 mS/cm, and 2 ppm O2. A line has been added to the fracture surface of the sampleannealed at 500°C for 45 min. to indicate the irradiated surface of the sample.

200 mm

1.0 dpa(1,1)

As. Irrad.

450°C/45 min.

500°C/45 min.

200 mm

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002 0.0025 0.003


t)1/2

, (cm)


CP-304 (1.0 dpa)CP-304 (2.5 dpa)

Jacobs 304 (~2.6 dpa) Katsura 316 (~11 dpa)Katsura 304 (~18 dpa)

200 mm200 mm

1.0 dpa(1,1)1.0 dpa(1,1)

As. Irrad.

450°C/45 min.

500°C/45 min.

200 mm200 mm

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002 0.0025 0.003


t)1/2

, (cm)


CP-304 (1.0 dpa)CP-304 (2.5 dpa)

Jacobs 304 (~2.6 dpa) Katsura 316 (~11 dpa)Katsura 304 (~18 dpa)


46

4.3 Discussion of Post-Irradiation Annealing Results

Simulations predicted that dislocation loops were removed preferentially over RIS.

Comparison of the measured and simulated annealing confirmed the accuracy of the simulations

for both segregation profiles and dislocation loops. The experimental separation of loops and

RIS via post-irradiation annealing is confirmed in the following section. The effects of Cr, Si,

and P segregation in IASCC are then examined by comparing the behavior of cracking and RIS

during annealing. In a similar manner, the importance of dislocation loops and hardness in

IASCC are determined.

4.3.1 Separation of RIS and Loops

The effects of post-irradiation annealing on microchemical changes were considerably

different than those on microstructure and hardening. Very little change was observed in grain

boundary composition or composition profiles in either alloy at most annealing conditions

examined. Even under the most extreme conditions (600°C for 90 min.) radiation-induced

segregation was largely unchanged. For all other conditions examined up to 600°C/90 min.,

measured microchemistry was virtually identical to the as-irradiated condition. Significant

changes in loop population and hardness were observed following post-irradiation annealing, in

contrast to the behavior observed for radiation-induced segregation. The data trends for the

removal of RIS (Figure 8) and dislocation microstructure (Figure 11) are superimposed over the

annealing of hardness shown in Figure 13. The removal of dislocation loops and hardness follow

the same trend and both are clearly removed preferentially to RIS. Indeed, annealing at 600°C

for 90 min. removed virtually all radiation-induced changes to yield stress. The preferential

removal of dislocation loops and hardening is consistent with the annealing simulations. Further,

RIS is not affected until 80% of the dislocation microstructure or hardening has been removed,

which is consistent with the simulation results shown in Figure 6 and supports that the

preferential removal of dislocation loops is partially due to competition for vacancies between

the removal of RIS and loops.

In summary, measured results from both this study and other studies indicate that

dislocation loops were removed preferentially over RIS. Comparison of the measured and


47

simulated annealing confirmed both the accuracy of the simulations and the preferential removal

of the dislocation loop microstructure during annealing. Since microchemical changes were

separated from microstructural changes and hardening, the cracking susceptibility following

annealing was compared to the as-irradiated condition. Although, the cracking susceptibility was

removed before RIS was truly isolated from the dislocation loops or hardening, careful

comparison of the annealing behavior of both cracking and RIS may help to help assess the

importance of RIS in IASCC. Direct comparisons of as-irradiated and annealed cracking

susceptibility to as-irradiated and annealed dislocation loop microstructure and hardening

provide further insight into the effect of each irradiation-induced change on IASCC.

4.3.2 RIS of Cr and Ni and IASCC

As shown in Figure 15 a.), the grain boundary Cr and Ni contents remain at the as-

irradiated level over the range of annealing conditions (plotted as a function of Fe-diffusion

distance) for both 1.0 and 2.5 dpa samples. However, IG cracking susceptibility clearly shows a

very rapid decrease with increasing annealing time or temperature and is completely eliminated

after annealing at 500°C for 45 min.

Given that the IG cracking susceptibility was removed before the grain boundary Cr

content begins to change indicates that Cr depletion alone cannot be a primary contributor to

IASCC. This result is consistent with the observed difference in cracking behavior of proton-

irradiated CP-304 and the companion CP-316 alloy studied previously [29]. At all doses, the CP-

316 exhibited more Cr depletion (in “w-shape” at 1.0 dpa or in “v-shape” at higher damage levels)

than the CP-304 alloy, yet the CP-316 alloy did not crack at doses up to 5.0 dpa. Jacobs [46] also

concluded in his study of neutron-irradiated 304 SS that the depletion of Cr does not appear to be a

primary causative factor in IASCC.”

Similarly, the grain boundary enrichment of Ni does not play an important role in the

mitigation of IASCC as cracking was removed during annealing with no measurable change in Ni

segregation. Grain boundary Ni content steadily increases with dose and is approximately twice

the bulk level by 5.0 dpa. However, cracking susceptibility also increases steadily with increasing

dose through 5.0 dpa, reinforcing that Ni enrichment cannot be a mitigating factor for IASCC.


48

Figure 15. Comparison of annealing behavior of major alloying elements, a.), minor alloyingelements, b.), dislocation loop line length, c.) and hardness, d.) with cracking susceptibility forCP-304 irradiated to 1.0 and 2.5 dpa.

4.3.3 RIS of Si and P and IASCC

Silicon and phosphorus enrichment and cracking susceptibility are compared in Figure 15

b.). During annealing, both Si and P enrichment are removed rapidly, at a rate similar to that of

cracking susceptibility. The similarities in recovery rates of both Si and P enrichment to that of

the cracking susceptibility potentially implicates both elements as contributors to IASCC.

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002


t)1/2 , (cm)

Cr Content%IG

1.0 dpa

2.5 dpa

a.)

Ni Content

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002


t)1/2 , (cm)

Si Content%IG

1.0 dpa

2.5 dpa

P Content

b.)

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002


t)1/2 , (cm)c.)

Sl

%IG

1.0 dpa

2.5 dpa

0

20

40

60

80

100

120

0 0.0005 0.001 0.0015 0.002


t)1/2 , (cm)d.)

%IG

1.0 dpa

2.5 dpa

Dsy


49

However, substantial literature data suggest that Si plays a beneficial role, if any at all

[45-50]. Chung [47] and Fukuya [48] both found that CP alloys with higher Si content were less

susceptible to IG cracking than high purity alloys irradiated to the same conditions. Tsukada

[49] reported that an alloy with high Si content failed at a higher strain than a high purity

reference alloy with a comparable fraction of IG cracking. Only Jacobs et al. [50] reported any

potential link between cracking and Si content. However, in a later study, Jacobs [46] reported

grain boundary Si enrichment had little impact on IASCC. Finally, the work of Cookson [45]

showed that increased Si content had a slightly beneficial effect on cracking of proton-irradiated

HP-304 SS. Combined, the experimental evidence from other studies shows that the role of Si

segregation in IASCC is minor.

The grain boundary P content of both 1.0 and 2.5 dpa samples, returned to the bulk

content following even the shortest anneal (400°C for 45 min.), faster than the mitigation of

cracking. While the percentage of IG cracking increases steadily with dose, the grain boundary P

content increased dramatically between 0.3 and 1.0 dpa. Thereafter, the P content remained

relatively unchanged until dropping back to the bulk level by 5.0 dpa (while cracking

susceptibility continued to increase). The trends during annealing and dose dependence reveal

that P enrichment during irradiation is the primary cause of IASCC, consistent with the work of

Cookson [45], Fukuya et al. [48], Tsukada et al. [49], Chung et al. [15], and Jacobs [51] who all

reported that P had either no effect or a slightly beneficial effect. While the results of this study

indicate that the effects of minor alloying elements such as Si and P on IASCC to be minor, other

minor alloying elements and impurities such as C, B, or N may be influential in IASCC.

4.3.4 Dislocation loops and IASCC

The total loop line length calculated from the dislocation loop microstructure is compared

to the cracking susceptibility in Figure 15 c.). During annealing, the cracking susceptibility was

mitigated before the total loop line length was significantly changed. No changes in loop density

or diameter were measured following annealing at 400°C or 450°C for 45 min. The lack of

change in loop line length after annealing at 500°C for 45 min. may be due to the inclusion of

unfaulted loops in the density count.


50

With irradiation, the total dislocation line length associated with dislocation loops

increased quickly with dose, reaching a saturation level between 1.0 and 3.0 dpa with little

increase between 3.0 and 5.0 dpa. The cracking susceptibility, however, continued to increase

considerably between 3.0 and 5.0 dpa. Further, dislocation densities and diameters measured in

the proton-irradiated CP-316 SS alloy were similar to those measured in the CP-304 alloy [29].

Despite the similar loop populations, a distinct difference in cracking susceptibility exists.

Therefore, the observed dislocation microstructure alone is not the primary mechanism for

IASCC.

4.3.5 Hardening and IASCC

The impact of radiation-induced hardening is assessed in Figure 15 d.), which plots the

change in yield stress and IG cracking susceptibility as a function of diffusion distance. During

annealing of both 1.0 and 2.5 dpa specimens, the degree of hardening remains at the as-irradiated

level for anneals at 400°C and 450°C. The hardening recovers slightly after annealing at 500°C

for 45 min., while cracking has been completely removed, indicating that cracking susceptibility

is not determined by hardness alone.

Like the dislocation loop microstructure, the change in yield stress increases steadily with

increasing radiation dose up to 3.0 dpa. Between 3.0 and 5.0 dpa, there is little increase in yield

stress change. Further, hardening of the proton-irradiated CP-316 alloy were similar to those

measured in the CP-304 of this study in spite of the difference in cracking behavior [29]. Thus,

radiation-hardening alone cannot explain the occurrence of IASCC in the alloys examined.

4.3.6 Other potential contributors to IASCC

The segregation of Cr, Ni, Si, and P were each found to be insufficient to cause IASCC

alone. Further, the dislocation microstructure and radiation-induced hardening alone did not

correlate with IASCC. All of these effects may contribute to IASCC, however, none seems to be

sufficient to cause IASCC alone. Therefore, some other feature or radiation-induced change

must be controlling the observed IASCC behavior. Possibilities include the segregation of other

minor elements such as C, B, or N, which are not typically measured, unresolved small defect


51

clusters, or a combination of effects.

4.3.7 Other elements

The CP-304 alloy contains minor elements such as C, N, and B, which were not analyzed

in this study. Quantitative analysis via STEM/EDS of B, C, or N is extremely difficult due to the

low x-ray yields of these light elements. Further, significant peak overlap with lower energy

peaks of major alloying elements such as Fe, Cr, and Ni makes even qualitative analysis difficult.

The CP-304 and CP-316 studied by Was et al. [29] have similar levels of C (0.16 and 0.18 at%,

respectively) and N (0.266 and 0.230 at%, respectively), making it difficult to explain the

difference in cracking susceptibility between the two alloys on the basis of these elements alone.

Boron has been identified as potentially having beneficial effects on IASCC [52]. Chung

et al. [15] reported that higher B concentration was beneficial in suppressing IASCC. In a study

by Jenssen et al. [16], alloys with low boron content showed higher susceptibility to IASCC than

those with higher bulk B levels. Boron was measured by Kenik et al. [53] using atom probe

analysis in both the unirradiated CP-304 and CP-316 alloys studied by Was [29]. The grain

boundary B content in the CP-316 SS was measured at 4.4 at%, considerably higher than the 1.4

at% content measured in the CP-304 SS, consistent with the observations of Chung [47] and

Jenssen [16] and supporting the potentially beneficial role of B. However, analysis of the grain

boundary B content must be performed on irradiated samples to confirm any potential role of B

in suppressing IASCC. Other analytical techniques such as atom probe might be useful,

however, techniques for creating atom-probe samples from proton-irradiated samples do not

currently exist.

4.3.8 Combination of RIS and microstructure

RIS of any one element, the formation of dislocation loops or increases in hardness

cannot alone account for the observed IASCC behavior. However, the presence of two or more

effects at a critical level may be responsible, as suggested by Chung et al. [15]. The results of this

study suggest that a combination of RIS, dislocation loops, and/or hardening controlling IASCC is

unlikely. Considering that the cracking susceptibility was mitigated with even the shortest


52

annealing treatment, prior to Cr or Ni segregation, dislocation loops, and hardening being altered,

no combination of the effects examined in this study can be solely responsible. However, changes

in both the dislocation loop population and hardness might occur at the same time and offset each

other, as discussed previously.

4.3.9 Small defect clusters

If none of the observed changes are responsible for the increase in cracking susceptibility

with increasing dose, some other irradiation-induced change may be responsible. Fine scale

radiation damage may be one of several possible contributors. During irradiation, interstitials or

vacancies survive the cascade event, either as individual defects or as small clusters. The small

clusters can take the form of vacancy or interstitial loops, stacking fault tetrahedra, di- or tri-

vacancy clusters, interstitial clusters or defect-impurity clusters. During deformation of the

sample, these small defects clusters will act to impede moving dislocations, similar to faulted

dislocation loops. Unfortunately, small defect structures are typically not characterized due to

the extreme difficulty in imaging this type of damage in TEM. Further, detailed and accurate

analysis of small defect clusters is complicated as routine sample preparation of TEM disks

typically involves the use of ion-milling to obtain an electron transparent region and can induce

small defect damage in metal samples. [54]

During post-irradiation annealing, small defects can be removed quickly via annihilation

due to their small size, or they can spontaneously dissociate, leaving individual interstitials or

vacancies behind. [55] Dislocation loops may absorb free interstitials to slightly increase the

mean radius or they may absorb vacancies and be removed, as depicted schematically in Fig. 16.

Detailed analysis of the smallest defects in the post-irradiation annealed specimens (CP-304 at

1.0 dpa) was performed with dark-field analysis. The bright-field, dark-field and dislocation

loop size distributions measured from the dark-field images are shown in Figure 10. Indeed, the

dislocation loop size distributions illustrated in Figure 10 do show that the smallest dislocation

loops were removed at the lowest annealing temperatures while the larger loops remained

unaltered. The increase in hardness due to an increase in loop size is offset by the removal of

the small defect clusters resulting in no observable change in loop density. Hence, the net effect

is a change in defect morphology with no discernible change in dislocation loop density or yield

strength calculated from measured hardness.


53

Figure 16. Schematic of annealing behavior of mixed population of black-dot damage anddislocations. During annealing, small defect clusters are removed or dissociate into individualinterstitials which may then be absorbed by dislocation loops. This process results in apopulation of larger loops with no change in density.

Further information about the nature of the small defect populations may be gained by

examining the post-irradiation annealing behavior of the samples irradiated at 360°C. In order to

simulate the effects of post-irradiation annealing for a population of interstitial defects, the post-

irradiation annealing model developed by Busby [56] was used. The model calculates the

changes in loop radius and density as a function of time at any given temperature for a

population of defects. The simulations were based on the methodology used by Burton [55] and

further details are described by Busby. [56]

The annealing model has been applied to examine the annealing behavior of the smallest

loops observed in the samples irradiated at 360°C and post-irradiation annealed. The loop

density measured in each size up to 6 nm is plotted in Figure 17 for both the samples annealed at

450°C and 500°C. The simulated distributions for the same annealing conditions are also plotted

in Figure 17 b.) for comparison. There is excellent agreement between the modeled and

measured loop size distributions. The modeled and measured dislocation density (m-2) for the

smallest loops (less than 6 nm) is plotted in Figure 17 c. Again, there is excellent agreement

between modeled and measured loop densities following annealing. This indicates that these

small defects in the samples irradiated at 360°C are, indeed, small, faulted dislocation loops.

This is in agreement with the recent results of Edwards et al. [57]

Defect clusters

Dislocation loops

Loops

Loop Growth

IASCC NO IASCC

Annealing

iii

Defect clusters

Dislocation loops

Loops

Loop Growth

IASCC NO IASCC

Annealing

iii


54

Figure 17. Comparison of measured and simulated annealing of small dislocation looppopulations from proton irradiations to 1.0 dpa at 360°C (plotted as Nd in c.)).

1020

1021

1022

1 2 3 4 5 6

Loop Diameter (nm)

500/45 minSimulated

1020

1021

1022

1 2 3 4 5 6

Loop Diameter (nm)

500/45 minMeasured

1020

1021

1022

1 2 3 4 5 6

Loop Diameter (nm)

450/45 minMeasured

1020

1021

1022

1 2 3 4 5 6

Loop Diameter (nm)

450/45 minSimulated

0

0.5

1

1.5

2

2.5

3

350 400 450 500 550

Annealing Temperature (oC)

Experimental results Simulation results

AsIrrad.

c.)

b.) a.)


55

4.3.10 Implications for IASCC

As noted earlier, the interaction between small loops and moving dislocations may result

in dislocation channeling. In the channeling process, the initial, moving dislocations encounter

dislocation loops, which act as barriers to the dislocations. To overcome these barriers,

dislocations annihilate or combine with the defects on the slip plane and continue to glide.

Subsequent dislocations will tend to glide along this same path, clearing out additional defects,

resulting in a channel that is free of defects. The elimination of the irradiation induced defects

along the primary slip planes results in localized deformation. These channels have been

observed in both pure metals [58] and in austenitic stainless steels following irradiation with

protons. [59] Channeling is more pronounced in alloys where the damage is in the form of a

high density of small defects as opposed to a dislocation network.

As noted by Bruemmer et al, [59] localized deformation can be detrimental to IASCC by

promoting dislocation pileups at grain boundaries. To accommodate local strain, a grain

boundary may absorb dislocations, emit new dislocations, create deformation microtwins, or

crack if no further strain can be accommodated. Further, dislocation channels may have a finite

life before pile-ups occur closing the channel to further dislocation glide. As available channels

for plasticity are eliminated, the grain boundary may not be able to accommodate further strain

by emitting new dislocations or microtwins. This is consistent with recent results from

Alexandreanu, [60] which examined IG cracking in Ni-base alloys. Alexandreanu found that

boundaries that accommodate deformation by absorbing dislocations are also more susceptible to

cracking.

Localized deformation has been observed in the post-irradiation annealed specimens of

this study. Micrographs of the slip bands observed in the samples annealed at 450°C and 500°C

for 45 minutes are compared in Figures 18 a and b. Slip bands were observed in both samples,

although fewer bands were observed in the sample annealed at 500°C for 45 minutes, which

failed with no evidence of IG cracking (25 slip band systems/mm2) while the sample annealed at

450°C for 45 minutes did have IG cracking and a higher density of bands (65/mm2). The CP-

304 sample irradiated to 0.3 dpa at 360°C also failed via ductile rupture, and the slip bands for

this condition are shown in Figure 18c. The number of slip bands in the 0.3 dpa sample

(22/mm2) is comparable to those observed in the sample annealed for 45 minutes at 500°C. If

fine-scale damage is a primary contributor to IASCC and removed rapidly during post-irradiation


56

annealing, a distinct difference in slip bands should be observed between annealed samples

which exhibited cracking (1.0 dpa sample annealed at 450°C for 45 minutes), and those which

failed via ductile rupture (1.0 dpa sample annealed at 500°C for 45 minutes). In the sample

annealed at 450°C for 45 minutes, the higher density of slip bands in samples exhibiting cracking

suggests that microstructural features such as fine scale clusters may be important in

deformation, and hence, cracking susceptibility.

Figure 18. Comparison of slip bands in CP 304 SS irradiated to 1.0 dpa at 360°C and annealedat a.) 450°C/45 minutes (IG with 65 slip band systems/mm2) or b.) 500°C/45 minutes (no IG and25 slip band systems/mm2). Also shown in c.) are the slip bands observed on a CP 304 SSsample irradiated to 0.3 dpa at 360°C (no IG with 22 slip band systems/mm2) and d.) 0.3 dpa at <75°C (no IG with 28 slip band systems/mm2). Samples strained at 3.5 x 10-7 s-1 in water at288°C, 0.2 mS/cm, and 2 ppm O2.

100 mmCP-304 0.3 dpaAs-Irrad.

Slipbands

100 mmCP-304 1.0 dpa500/45 min100 mm

CP-304 1.0 dpa450/45 min

Slipbands Slip

bands

No IG

IG No IG

c.)

a.) b.)

50 mmCP-304 0.3 dpaLT As-irrad

Slipbands

d.)

100 mmCP-304 0.3 dpaAs-Irrad.

Slipbands

100 mmCP-304 1.0 dpa500/45 min100 mm

CP-304 1.0 dpa450/45 min

Slipbands Slip

bands

No IG

IG No IG

c.)

a.) b.)

50 mmCP-304 0.3 dpaLT As-irrad

Slipbands

d.)


57

Finally, note the mitigation of cracking during post-irradiation annealing shown in Figure

14 for this study is very similar to that of Jacobs and Katsura (with the exception of Katsura’s

304 SS at a diffusion distance of 0.0016 cm). The similarity in cracking mitigation, despite the

wide disparity in as-irradiated dose, indicates that cracking is controlled by a feature or

condition, which has reached a critical level by 1.0 dpa (since cracking was observed at this

dose). The fine defect damage reaches its saturation level very quickly during irradiation. [61]

The rapid removal of a dense population of small obstacles impeding dislocation motion could

result in a change in deformation mode and, hence, cracking behavior during annealing. This is

consistent with the work of Bailat et al., [62] who reported a possible correlation between

deformation mode and IASCC in neutron-irradiated stainless steels. Bailat et al. found localized

deformation (channeling and microtwinning) in a susceptible 304 SS alloy using TEM analysis.

A companion 316 alloy irradiated to the same conditions did not exhibit the same characteristics

of localized deformation and was resistant to IASCC.

Low temperature irradiation is another means of examining fine scale radiation damage.

Low temperature irradiations can create a relevant irradiation-induced microstructure, consisting

primarily of a high density of small defect clusters, or black dot damage without the

accompanying RIS. The following section presents results of the low temperature irradiations

designed to isolate the irradiated microstructure.

4.4 Isolating the Dislocation Microstructure

Low-temperature irradiations were conducted to create a high density of small defect

clusters. Two irradiation programs were completed, the first to a dose of 0.3 dpa at T<75°C, and

the second to doses of 0.5 to 1.5 dpa in varying combinations of low (<75°C) and high (360°C)

temperature irradiation designed to achieve a stable irradiated microstructure. The increase in

hardness of samples irradiated to 0.3 dpa at T<75°C is about 75 kg/mm2, which is higher than the

value of 65 kg/mm2 obtained after 1.0 dpa at 360°C. The high degree of hardening in the

specimens irradiated to 0.3 dpa at T<75°C indicates that a very high density of defects has

indeed been created. Subsequent annealing at temperatures between 350°C and 500°C resulted

in annealing of the defects and a reduction in hardness, table 7 and Fig. 19. However, no visible

defects have been observed in any of the as-irradiated or post-irradiation annealed specimens.

TEM has the capability to resolve dislocation loops down to about 1.0 nm in diameter, and to


58

detect smaller loops. This indicates that the defects created during low temperature irradiation

are less than 1 nm in diameter.

Table 7. Hardness of CP304 SS following irradiation and annealing.

Condition DHv (kg/mm2) Relative Hardness

Reference condition: 1.0dpa@360°C

60 1.0

As-irrad. 0.3 dpa @ 50°C 85 1.42

As-irrad. + 3.8 hr @ 350°C 66 1.09

As-irrad. + 0.5 hr @ 400°C 35 0.59

As-irrad. + 0.5 hr @ 450°C 27 0.45

As-irrad. + 0.5 hr @ 500°C 23 0.36

Figure 19. Change in hardness of CP304 SS following annealing of samples irradiated to 0.3dpa at T<75°C.

0.0040.0030.0020.0010.00020

40

60

80

100

1/T (K-1)

Har

dnes

s, D

Hv (

kg/m

m2 )

350°C:3.8 hr

As-irrad.


59

Figure 20 confirms that following low temperature irradiation to 0.3 dpa, there is no

evidence of RIS. As shown, chromium, nickel and iron contents are nearly flat across the grain

boundary. Even silicon, which segregates strongly under irradiation, shows no sign of RIS,

indicating that irradiation at low temperature was indeed successful in suppressing RIS.

Figure 20. Composition profiles for Cr, Ni, Fe and Si across grain boundaries of CP304 SSfollowing low temperature irradiation at T<75°C showing the lack of RIS.

A second series of irradiations was conducted to try to stabilize and grow the defect

clusters using combinations of low and high temperature irradiations. Figure 21 shows the

scheme for combinations of low and high temperature irradiations. Both TEM bars and SCC

bars were irradiated to doses of 0.5 dpa on each step. Consequently, samples received total

doses of between 1.0 and 1.5 dpa. Hardness results are given in table 8 and confirm that

irradiation at low temperature or at high temperature followed by low temperature result in

greater hardening than at high temperature or low temperature followed by high temperature.

The results are quite consistent and indicate again that the density of very small defects

following low temperature irradiation must be very high.

Stress corrosion cracking tests in normal water chemistry were conducted on the

combination of low and high temperature irradiations shown in table 9. The stress-strain curves

are given in Fig. 22. Note that with one exception, the samples all reached similar levels of

stress and strain. Post-test inspection in the SEM revealed no IGSCC in these samples.

0

5

10

15

20

25

-12 -8 -4 0 4 8 12


Protons at T<75°C0.3 dpa

Cr

Ni

60

62

64

66

68

70

72

74

0

0.4

0.8

1.2

-12 -8 -4 0 4 8 12


Protons at T<75°C0.3 dpa

Fe

Si


60

Figure 21. Scheme for achieving a stable irradiated microstructure by combining hightemperature (T=360°C) and low temperature (T<75°C) irradiation in increments of 0.25 dpa.

Table 8. Hardness before and after various irradiations to 0.5 dpa at either T<75°C or at 360°C.All hardness units are in kg/mm2.Sample # Unirradiated

hardness

Hardness

at 0.5

dpa@

360°C

DH from

previous

condition

Hardness

at 0.5 @

75°C

DH from

previous

condition

Hardness

at 0.5 dpa

@ 360°C

DH from

previous

condition

T2 203 258 55 - - 258 0

T3 206 261 55 - - 258 -3

T4 207 256 49 313 57 - -

T5 205 - - 283 78 - -

T6 205 - - 298 93 248 -50

S4T4 T5 S5S1 S2 S3 S6T2 T3 T6 T7T1 S4 T5

KEY:

H - H H - LLL - H H - L - H


61

This is not surprising in retrospect as the low temperature irradiated microstructure is not stable

at the SCC test temperatures. Figure 23 shows the decrease in hardness following an anneal for

various times at temperatures at the SCC test temperature of 288°C. The softening in the CERT

test is likely to be even more extensive than shown here as these tests ran for a duration of about

15 days.

Table 9. Results of constant extension rate test in 288°C BWR normal water chemistry.

Sample

IrradiationH

istory

% Strain at

failure

UT

S(ksi)

# of cracksin irradiatedregion

Total C

rackL

ength onirradiatedSurface(_m

)

Location of

Failure

FractureM

ode

% IG

% T

G

% D

uc.

S2 1.0 dpa360-360

36.2 63.6 1 187052

IrradiatedRegion

IG+TG+Ductile

2.0 38.8 59.7

S3 1.5 dpa360-75-360

35.8 63.9 4 42812

Threads NA NA NA NA

S4 1.0 dpa360-75

31.6 63.0 1 190060

IrradiatedRegion

TG+Ductile

0 18 82

S6 1.0 dpa 75-360

31.8 55.9 3 1314

Unirr.Region

NA NA NA NA

Ref. 1.0 dpa 360 23.6 64.9 1 142060

IrradiatedRegion

IG+TG+Ductile

2.2 61.7 36.1

While the current TEM analysis techniques have failed to reveal any information about

the nature of these defects, some insight may be gained using the hardness results and

simulations of post-irradiation annealing. Using the measured change in hardness for the as-

irradiated specimen, the corresponding change in yield stress can be calculated using Eq. (1). An

expected change in yield stress can also be calculated from irradiated microstructure using the

dispersed barrier-hardening model [36]:

(7)

where a is the barrier strength (0.2 for small clusters [63], M is the Taylor factor (3.06), m is the

shear modulus (76 GPa), b is the Burgers vector (0.255 nm), N is the defect density and d is the

defect diameter. Using the calculated change in yield stress of 263 MPa, determined from the

Dsy = aMmb(Nd)1/2


62

hardness measurements, Nd can be determined as 5 x 1014 m-2. If a mean cluster size of 0.5 nm

is assumed (being approximately half of the smallest loop size measured in the samples

irradiated at 360°C), the defect density is approximately 1 x 1024 m-3. While this density is

considerably higher than the loop density measured in the same alloy irradiated at 360°C, it is

not an unreasonable density for small defect clusters created during low temperature irradiation.

There is little data on microstructure development at low temperatures, although in a review

paper, Zinkle [64] notes that the saturation density of small defect clusters may be greater than

1024 m-3, and that this density is attained after doses of about 0.1 dpa. Further, the data from

Higgy and Hammad [65] suggest that mechanical property measurements on 304 and 316 SS

after neutron irradiation at <100°C suggest that defect cluster saturation occurs before 0.1 dpa,

consistent with the results of this study.

Figure 22. Stress-strain curves of CP304SS samples tested in normal water chemistry consistingof 288°C water containing 2 ppm O2 and conductivity of 0.2 mS/cm. Samples were irradiated tothe following conditions: a) 0.25 dpa @ T<75°C + 0.25 dpa @ 360°C, b) 0.25 dpa @ 360°C +0.25 dpa @ 360°C, c) 0.25 dpa @ 360°C + 0.25 dpa @ T<75°C, and d) 0.25 dpa @ 360°C +0.25 dpa @ T<75°C + 0.25 dpa @ 360°C.

0

40

80

0 5 10 15 20 25 30 35 40

Stre

ss (k

si)

Strain (%)

Low-HighSTF: 31.84%Max Stress: 55.9 ksi

0

40

80

0 5 10 15 20 25 30 35 40

Stre

ss (k

si)

Strain (%)

High-HighSTF: 36.18%Max Stress: 63.6 ksi

0

40

80

0 5 10 15 20 25 30 35 40

Stre

ss (k

si)

Strain (%)

High-LowSTF: 31.55%Max Stress: 63.0 ksi

0

40

80

0 5 10 15 20 25 30 35 40

Stre

ss (k

si)

Strain (%)

High-Low-HighSTF: 35.8%Max Stress: 63.9 ksi


63

Figure 23. Change in hardness of CP304SS following irradiation to 0.3 dpa and annealing at288°C for various times up to 24 hours.

The model developed by Busby has also been used to simulate the annealing of samples

irradiated at low temperature. The as-irradiated defect density determined earlier was used as the

initial condition in the simulations. A mean diameter of 0.5 nm, as discussed above, was

assumed with the defect population being uniformly distributed between 0 to 1 nm. The

experimental data is plotted along with simulation results in Figure 24. All results are plotted as

Nd (calculated from hardness for the experimental results) as a function of annealing time.

Clearly, the simulation at 350°C (solid line) predicts that the defect population is removed much

slower than experimentally observed. Indeed, a temperature of 475°C is required for the

simulations to match the experimental results. This discrepancy can be explained as follows.

Change in Hardness @ 288C

0

50

100

150

200

250

300

350

Un-Irr As Irr 12Hr @288C 24Hr @288C

Vic

kers

Har

dn

ess

Series1


64

First, the defect clusters in the simulations can only shrink by vacancy absorption, as there is no

mechanism in the model for spontaneous dissolution. More significantly, the model developed

by Busby only considers clusters that are interstitial in nature (as most dislocation loops are

considered to be). Fine scale clusters are likely a mixed population of interstitial and vacancy

type. Small vacancy clusters created under low temperature irradiation conditions are likely to

dissociate rapidly following even low temperature annealing. Further, if vacancy and/or

interstitial defect clusters spontaneously dissociate, this will provide additional vacancies and/or

interstitials, which may be absorbed at other clusters. All these factors combined would result in

accelerated annealing in the simulations, similar to what is experimentally observed.

Figure 24. Comparison of annealing of samples irradiated at low temperatures with modelsimulations of annealing of small defect clusters.

Further simulations were performed using a model developed by Simonen. [66] The

model developed by Simonen accounts for both interstitial and vacancy type loops or clusters.

Clusters can absorb or emit vacancies or interstitials resulting in growth or reduction,

1013

1014

1015

0 5 10 15 20 25 30

Annealing Time (minutes)

350oC

400oC

450oC

475oC

Experimental data: anneals at 350oC

(from hardness measurements)

Simulated annealing at various temperatures


65

respectively, for a vacancy cluster (and vice-versa for an interstitial cluster). The simulation

predictions for hardness after annealing are plotted in Fig. 25 as a function of adjusted time at

temperature (a simple variable which combines accounts for changes in both annealing time and

temperature). The hardness data from previous studies of post-irradiation annealing on proton-

irradiated samples are also plotted for comparison. The model and measured results from higher

temperature irradiations are in excellent agreement. The results from the low temperature

irradiations and anneals are also plotted in Fig. 25. However, these measured recovery occurs

much faster than that predicted by Simonen’s model. The discrepancy between modeled and

measured results indicates that the assumed ratio of vacancy to interstitial clusters does not

represent the low temperature irradiations. In this simulation, 40% of the clusters were assumed

to be vacancy-type. However, the rapid annealing of samples irradiated at low temperature

indicate that this percentage is likely to be 50% or higher. Nevertheless, agreement between

model and experiment is excellent.

Figure 25. Comparison of annealing of samples irradiated at low temperatures with predictionsfrom model developed by Simonen.

-0.2

0.0

0.2

0.4

0.6

0.8

1.0

10-20 10-18 10-16 10-14 10-12 10-10

400 C

450 C

500 C

550 C

600 C

650 C

Calculated

Adjusted Time at Temperature, s

Proton Irradiation at

Low TemperatureIrradiations

360oC to 1.0 dpa


66

2.0 CONCLUSIONS

A radiation annealing model was developed that describes the elimination of dislocation

loops by vacancy absorption. Results showed that there were indeed, time-temperature

annealing combinations that leave the radiation induced segregation profile largely unaltered

while the dislocation microstructure is significantly reduced. Proton irradiation of 304 stainless

steel irradiated with 3.2 MeV protons to 1.0 or 2.5 dpa resulted in grain boundary depletion of

chromium and enrichment of nickel and a radiation damaged microstructure. Post irradiation

annealing at temperatures of 500 – 600°C for times of up to 45 min. removed the dislocation

microstructure to a greater degree with increasing temperatures, or times at temperature, while

leaving the radiation induced segregation profile relatively unaltered. Constant extension rate

tensile (CERT) experiments in 288°C water containing 2 ppm O2 and with a conductivity of 0.2

mS/cm and at a strain rate of 3 x 10-7 s-1 showed that the IASCC susceptibility, as measured by

the crack length per unit strain, decreased with very short anneals and was almost completely

removed by an anneal at 500°C for 45 min. This annealing treatment removed about 15% of the

dislocation microstructure and the irradiation hardening, but did not affect the grain boundary

chromium depletion or nickel segregation, nor did it affect the grain boundary content of other

minor impurities. These results indicate that RIS is not the sole controlling feature of IASCC in

irradiated stainless steels in normal water chemistry.

The isolation of the irradiated microstructure was approached using low temperature

irradiation or combinations of low and high temperature irradiations to achieve a stable,

irradiated microstructure without RIS. Experiments were successful in achieving a high degree

of irradiation hardening without any evidence of RIS of either major or minor elements. The low

temperature irradiations to doses up to 0.3 dpa at T<75°C were also very successful in producing

hardening to levels considerably above that for irradiations conducted under nominal conditions

of 1 dpa at 360°C. However, the microstructure consisted of an extremely fine dispersion of

defect clusters of sizes that are not resolvable by either transmission electron microscopy (TEM)

or small angle x-ray scattering (SAXS). The microstructure was not stable at the 288°C IASCC

test temperature and resulted in rapid reduction of hardening and presumably, annealing of the

defect clusters at this temperature as well. Nevertheless, the annealing studies showed that


67

treatments that resulted in significant decreases in the hardening produced small changes in the

dislocation microstructure that were confined to the elimination of the finest of loops (~1 nm).

These results substantiate the importance of the very fine defect microstructure in the IASCC

process.

The results of this program provide the first definitive evidence that RIS is not the sole

controlling factor in the irradiation assisted stress corrosion cracking of austenitic stainless steels

in normal water chemistry. Earlier research has suggested that RIS may not play the dominant

role it has recently been afforded, but the results of this program are the first to definitively

establish this role. The program has also shown that the fine defect structure is implicated in the

IASCC process and likely plays a role in the observed cracking of core components. The results

of this project provide the basis and the motivation for further investigation into the mechanism

of IASCC in which fine defect clusters may play a crucial, if not a defining role.


68

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54. D. J. Barber, Ultramicroscopy, 52 (1993) p. 101-125.55. B. Burton, Materials Science and Technology, Vol. 8, 1992 pp. 602.56. J. T., Busby, G. S. Was, and E. A Kenik, J. Nucl. Mater., Vol. 302, 2002, pp. 20-40.57. D. J. Edwards, E. P. Simonen, and S. M. Bruemmer, J. Nucl. Mater., 317 (2003), p. 13-31.58. P. J. Maziasz, and C. J. McHargue, International Metal Review., Vol. 32, 1987, pp. 190.59. S. M. Bruemmer, J. I. Cole, J. L. Brimhall, R. D. Carter, and G. S. Was, Proceedings of the

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published in the Phil. Mag. A., 2003.


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7.0 PUBLICATIONS

J. T. Busby, M. M. Sowa, G. S. Was and E. P. Simonen, “Post-irradiation Annealing of SmallDefect Clusters,” submitted to Phil. Mag. A.

E. P. Simonen, D. J. Edwards, B. W. Airey, S. M. Bruemmer, J. T. Busby and G. S. Was,“Annealing Stages in Neutron-Irradiated Austenitic Stainless Steels,” submitted to Phil. Mag. A.

G. S. Was, J. Busby, M. Sowa, M. Hash and R. Dropek, “Role of Irradiated Microstructure andMicrochemistry in Irradiation Assisted Stress Corrosion Cracking,” submitted to Phil. Mag. A.

G. S. Was and J. T. Busby, “Recent Developments in Irradiation Assisted Stress CorrosionCracking,” 101h Int’l Conf. Environmental Degradation of Materials in Nuclear Power Systems –Water Reactors, American Nuclear Society, LaGrange Park, IL, invited.

J. T. Busby, G. S. Was and E. A. Kenik, “Role of Radiation-Induced Segregation in IASCC ofAustenitic Stainless Steels,” J. Nucl. Mater., 302 (2002) 20-40.

G. S. Was, J. T. Busby, T. Allen, E. A. Kenik, A. Jenssen, S. M. Bruemmer, J. Gan, A. D.Edwards, P. Scott and P. L. Andresen, “Emulation of Neutron Irradiation Effects with Protons:Validation of Principle,” J. Nucl. Mater., 300 (2002) 198-216.

J. T. Busby, M. M. Sowa and G. S. Was, “The Role of Fine Defect Clusters in Irradiation-Assisted Stress Corrosion Cracking of Proton-Irradiated 304 Stainless Steel,” Effects ofRadiation on Materials: 21st International Symposium, ASTM STP, M. L. Grossbeck, T. R.Allen, R. G. Lott and A. S. Kumar, Eds., American Society for Testing and Materials, WestConshohocken, PA, 2002.

J. T. Busby, G. S. Was and E. A. Kenik, “Isolation of the Role of Radiation-Induced Segregationin Irradiation Assisted Stress Corrosion Cracking of Proton-Irradiated Austenitic StainlessSteels,” 10th Int’l Conf. Environmental Degradation of Materials in Nuclear Power Systems –Water Reactors, NACE International, Houston, TX, 2002.

J. T. Busby, J. Gan, M. Daniels, G. S. Was, S. M. Bruemmer, D. J. Edwards and E. A. Kenik,“Microchemistry and Microstructure Evolution in Proton-Irradiated Austenitic Stainless Steels,”Proc. 9th International Conference on Environmental Degradation of Materials in Nuclear PowerSystems - Water Reactors, Minerals, Metals and Materials Society, Warrendale, PA, 1999, pp.1089-1098.


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8.0 STUDENTS

A total of three students supported on this program received graduate degrees from theDepartment of Nuclear Engineering, University of Michigan.

Jeremy Busby received his PhD thesis on “Role of Radiation Induced Segregation in IrradiationAssisted Stress Corrosion Cracking,” in December, 2000.

Ben Grambau received his MS degree in December, 2001.

Shawn Bilek received her MS degree in August, 2001.

Mechanism of Irradiation Assisted Cracking of Core ... · Table 3: Summary of grain boundary composition measurements* on post-irradiation annealed CP-304 (irradiated to1.0 and 2.5

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